Nanoparticles
Edited by Günter Schmid
Nanoparticles: From Theory to Application. Edited by Günter Schmid
Copyright 8 2004 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
ISBN: 3-527-30507-6
Further Titles of Interest
M. K€
ohler, W. Fritzsche
Nanotechnology
2004, ISBN 3-527-30750-8
C. N. R. Rao, A. M€
uller, A. K. Cheetham (Eds.)
The Chemistry of Nanomaterials
Synthesis, Properties and Applications
2003, ISBN 3-527-30686-2
P. M. Ajayan, L. S. Schadler, P. V. Braun
Nanocomposite Science and Technology
2003, ISBN 3-527-30359-6
F. Caruso (Ed.)
Colloids and Colloid Assemblies
2003, ISBN 3-527-30660-9
Z. L. Wang (Ed.)
Characterization of Nanophase Materials
1999, ISBN 3-527-29837-1
J. H. Fendler (Ed.)
Nanoparticles and Nanostructured Films
Preparation, Characterization and Applications
1998, ISBN 3-527-29443-0
Nanoparticles
From Theory to Application
Edited by Günter Schmid
Prof. Dr. Günter Schmid
Institute of Inorganic Chemistry
University Duisburg-Essen
Universitätsstr. 5-7
45117 Essen
Germany
1:09 am, 4/22/05
9 This book was carefully produced. Nevertheless, authors, editor and publisher
do not warrant the information contained
therein to be free of errors. Readers are
advised to keep in mind that statements,
data, illustrations, procedural details or
other items may inadvertently be
inaccurate.
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from the British Library.
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publication in the Deutsche
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( 2004 WILEY-VCH Verlag GmbH & Co.
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All rights reserved (including those of
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ISBN 3-527-30507-6
v
Contents
1
General Introduction
Günter Schmid
2
Quantum Dots 4
Wolfgang Johann Parak, Liberato Manna, Friedrich Christian Simmel,
Daniele Gerion, and Paul Alivisatos
Introduction and Outline 4
Nanoscale Materials and Quantum Mechanics 6
2.1
2.2
2.2.1
2.2.2
2.3
2.4
2.4.1
2.4.2
2.4.3
2.4.4
2.5
2.6
2.6.1
2.6.2
2.6.3
2.7
2.7.1
2.7.2
2.7.3
2.7.4
2.8
2.8.1
2.8.2
2.8.3
2.8.4
1
Nanoscale Materials as Intermediate between Atomic and Bulk
Matter 6
Quantum Mechanics 7
From Atoms to Molecules and Quantum Dots 8
Shrinking Bulk Material to a Quantum Dot 10
Three-Dimensional Systems (Bulk Material) 10
Two-Dimensional Systems 14
One-Dimensional Systems (Quantum Wires) 18
Zero-Dimensional Systems (Quantum Dots) 20
Energy Levels of a (Semiconductor) Quantum Dot 21
Varieties of Quantum Dots 24
Lithographically Defined Quantum Dots 25
Epitaxially Self-Assembled Quantum Dots 27
Colloidal Quantum Dots 28
Optical Properties of Quantum Dots 29
Absorption and Emission Spectra 29
Spectral Diffusion and Blinking 31
Metal Nanoparticles 33
Overview of Some Selected Applications 33
Some (Electrical) Transport Properties of Quantum Dots 34
Coulomb Blockade: Basic Theory and Historical Sketch 35
Single-Electron Tunneling 37
Tunneling Transport – Line Shape of Conductance Peaks 41
Some Applications 42
References 44
vi
Contents
3
Syntheses and Characterizations
3.1
3.1.1
Semiconductor Nanoparticles 50
Synthesis and Characterization of II–VI Nanoparticles
50
50
Alexander Eychmüller
3.1.1.1 Historical Review 50
3.1.1.2 Thiol-Stabilized Nanoparticles 57
3.1.1.3 The ‘‘TOP/TOPO’’ Synthesis 61
3.1.1.4 Core-Shell Nanocrystals 63
3.1.1.5 Quantum Dot Quantum Wells 65
References 71
3.1.2 Synthesis and Characterization of III–V Semiconductor
Nanoparticles 79
Uri Banin
3.1.2.1 Introduction 79
3.1.2.2 Synthetic Strategy 81
3.1.2.3 InAs and InP Nanocrystals 84
3.1.2.4 III–V Core/Shell Nanocrystals – Synthesis and Characterization
Acknowledgement 105
References 105
3.1.3 Synthesis and Characterization of Ib-VI Nanoclusters 107
89
Stefanie Dehnen, Andreas Eichhöfer, John F. Corrigan, and Dieter Fenske
3.1.3.1 Introduction 107
3.1.3.2 Chalcogen-Bridged Copper Clusters 108
3.1.3.3 Chalcogen-Bridged Silver Clusters 158
3.1.3.4 Selenium-Bridged Gold Clusters 181
References 183
3.2
Synthesis of Metal Nanoparticles 185
3.2.1 Noble Metal Nanoparticles 186
John S. Bradley and Günter Schmid
3.2.1.1 Introduction 186
3.2.1.2 History and Background 186
3.2.1.3 Generation of Metal Particles 187
3.2.1.4 Stabilization of Colloidal Metal Particles in Liquids 189
3.2.1.5 Synthetic Methods for the Preparation of Colloidal Transition
Metals 191
3.2.2 Synthesis and Characterization of Magnetic Nanoparticles 199
Dmitri V. Talapin, Elena V. Shevchenko, Horst Weller
3.2.2.1 Introduction 199
3.2.2.2 Organometallic Synthesis of Magnetic Nanoparticles 202
3.2.2.3 Size- and Shape-Dependent Magnetic Properties of Ferromagnetic Metal
Nanoparticles 220
References 230
4
Organization of Nanoparticles
4.1
Semiconductor Nanoparticles
Alexander Eychmüller
239
239
Contents
4.1.1
4.1.2
4.1.3
4.1.4
4.2
Molecular Crystals 239
Superlattices of Semiconductor Nanoparticles 242
Layers of Semiconductor Nanocrystals 244
Coupling of Semiconductor Nanocrystals 249
Metal Nanoparticles 251
Günter Schmid, Dmitri V. Talapin, and Elena V. Shevchenko
4.2.1 Three-dimensional Organization of Metal Nanoparticles
4.2.2 Ordered Monolayers of Metal Nanoparticles 260
4.2.2.1 Self-Assembly 261
4.2.2.2 Guided Self-Assembly 268
4.2.2.3 Aimed Structures 277
4.2.3 One-dimensional Arrangements 280
References 298
251
5
Properties
5.1
5.1.1
Semiconductor Nanoparticles 305
Optical and Electronic Properties of III–V and II–VI Nanoparticles
305
305
Uri Banin and Oded Millo
5.1.1.1 Introduction 305
5.1.1.2 Semiconductor Nanocrystals as Artificial Atoms 306
5.1.1.3 Theoretical Descriptions of the Electronic Structure 313
5.1.1.4 Atomic-like States in Core/Shell Nanocrystals – Spectroscopy and
Imaging 314
5.1.1.5 Level Structure of CdSe Quantum Rods 319
5.1.1.6 Optical Gain and Lasing in Semiconductor Nanocrystals 320
Acknowledgements 322
5.1.2 Optical and Thermal Properties of Ib–VI Nanoparticles 322
Stefanie Dehnen, Andreas Eichhöfer, John F. Corrigan, and Dieter Fenske
5.1.2.1 Optical Spectra and Thermal Behavior 322
5.1.2.2 Optical Spectra in the Solid State 326
5.2
Electrical Properties of Metal Nanoparticles 328
Ulrich Simon
5.2.1 Introduction 328
5.2.2 Physical Background and Quantum Size Effect 329
5.2.3 Single-Electron Tunneling 331
5.2.3.1 Working Principles of Metal-Based Single-Electron Tunneling
Devices 333
5.2.4 Single-Electron Tunneling in Nanoparticles 338
5.2.4.1 Thin Film Structures and STM Single-Electron Systems 340
5.2.4.2 SET on Single Chemically Tailored Nanoparticles 342
5.2.4.3 Metal Nanoparticles in Single-Electron Devices 345
5.2.4.4 Mesoscopic Arrays with Nanoelectrodes 347
5.2.4.5 Chemical Switching and Gating of Current through Nanoparticles 352
5.2.5 From Single-Particle Properties to Collective Charge Transport 353
5.2.5.1 One-Dimensional Arrangements 356
5.2.5.2 Two-Dimensional Arrangements 358
vii
viii
Contents
5.2.5.3 Three-Dimensional Arrangements
References 362
6
6.1
6.2
6.2.1
6.2.2
6.2.3
6.2.4
6.3
6.3.1
6.3.2
6.4
6.4.1
6.4.2
6.4.3
6.5
6.5.1
6.5.2
6.6
6.6.1
6.6.2
6.6.3
6.7
7
359
Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and
Applications 368
Eugenii Katz, Andrew N. Shipway, and Itamar Willner
Introduction 368
The Synthesis and Properties of Biomaterial-Functionalized
Nanoparticles 371
Functionalization by Electrostatic Adsorption 371
Functionalization by Chemisorption of Thiol-Derivatized
Biomaterials 372
Functionalization by Specific Interactions 373
Properties of Nanoparticle-Biomaterial Composites 374
Biomaterial-Functionalized Nanoparticles for Controlled Chemical
Reactivity 376
Controlling DNA Reactivity 376
Reporting on DNA Reactivity 377
The Aggregation of Biomaterial-Functionalized Nanoparticles 377
Receptor-Induced Aggregation of Guest-Functionalized
Nanoparticles 379
Nucleic Acid-Functionalized Nanoparticles for Controlled
Aggregation 382
Composite Assemblies of Nucleic Acids, Proteins, and
Nanoparticles 385
Assembly of Biomaterial-Nanoparticle Architectures on Surfaces 390
Assembly of Layered Nanoparticle-Protein Assemblies 391
Nucleic Acid – Nanoparticle Architectures on Surfaces 392
Functional Biomaterial-Nanoparticle Structures on Surfaces for Sensoric
and Electronic Applications 397
Nanoparticle-Biomaterial Conjugates for Optical Sensing and
Analysis 398
Electronic Transduction of Biosensing Events by Nanoparticle-BiomaterialFunctionalized Arrays 400
Biomaterial-Nanoparticle Arrays for Electronic Circuitry 407
Biomaterial-Functionalized Magnetic Particles 409
Acknowledgment 415
References 415
Conclusions and Perspectives
All Authors
422
ix
Contributors
Prof. Paul Alivisatos
University of California, Berkeley
Department of Chemistry
Berkeley, CA 94720-1460
USA
Dr. Uri Banin
Department of Physical Chemistry
The Hebrew University
Jerusalem 91904
Israel
Prof. John S. Bradley
Dept. of Inorganie and Materials Chemistry
School of Physical Sciences
University of Hull
Cottingham Road
Hull HU6 7RX
UK
Prof. John F. Corrigan
Department of Chemistry
The University of Western Ontario
Chemistry Building
London, Ontario, N6A 5B7
Canada
Dr. Alexander Eychmüller
Institut für Physikalische Chemie
Universität Hamburg
Bundesstraße 45
20146 Hamburg
Germany
Prof. Dr. D. Fenske
Universität Karlsruhe
Institut für Anorganische Chemie
Engesserstraße
76128 Karlsruhe
Germany
Daniele Gerion
University of California, Berkeley
Department of Chemistry
Berkeley, CA 94720-1460
USA
Prof. Eugenii Katz
Department of Chemistry
The Hebrew University of Jerusalem Givat
Ram
91904 Jerusalem
Israel
Dr. Stefanie Dehnen
Institut für Anorganische Chemie
Universität Karlsruhe
Kaiserstraße 12
76131 Karlsruhe
Germany
Dr. Liberato Manna
National Nanotechnology Lab of INFM
Via Arnesano
73100 Lecce
Italy
Dr. Andreas Eichhöfer
Forschungszentrum Karlsruhe
Institut für Nanotechnologie
Postfach 3640
76021 Karlsruhe
Germany
Dr. W. J. Parak
Universität München
Lehrstuhl für Angewandte Physik
Amalienstr. 54
80799 München
Germany
Nanoparticles: From Theory to Application. Edited by Günter Schmid
Copyright 8 2004 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
ISBN: 3-527-30507-6
x
Contributors
Prof. Dr. Günter Schmid
Institut für Anorganische Chemie
Universitätsstraße 5-7
Universität Duisburg-Essen
45117 Essen
Germany
Prof. Dr. Ulrich Simon
RWTH Aachen
Anorganische Chemie
Professor Pirlet Straße 1
52054 Aachen
Germany
Dr. Elena V. Shevchenko
Institute of Physical Chemistry
University of Hamburg
Bundesstraße 45
20146 Hamburg
Germany
Dr. Dmitri V. Talapin
Institute of Physical Chemistry
University of Hamburg
Bundesstraße 45
20146 Hamburg
Germany
Dr. Andrew N. Shipway
Department of Chemistry
The Hebrew University of Jerusalem Givat
Ram
91904 Jerusalem
Israel
Prof. Horst Weller
Institute of Physical Chemistry
University of Hamburg
Bundesstraße 45
20146 Hamburg
Germany
Dr. Friedrich C. Simmel
Sektion Physik der LMU
Geschwister-Scholl-Platz 1
80539 München
Germany
Prof. Italmar Willner
The Institute of Chemistry
The Hebrew University of Jerusalem
91904 Jerusalem
Israel
1
1
General Introduction
Günter Schmid
In 1994 the book entitled ‘‘Clusters and Colloids – from Theory to Applications’’
was published with the goal of summarizing the state of the art in the field of
metal clusters and colloids. Nine years ago, interest was mainly focused on synthetic and structural aspects. Theoretical considerations were limited to small
molecule-like organometallic clusters. The very first hints of important novel electronic properties of nanosized particles stimulated further activity. Practical applications could be foreseen – at least if the will was there.
Nine years of worldwide revolutionary developments in nanoscience, combining
physics, chemistry, material science, theory and even biosciences, have brought us
to another level of understanding. ‘‘Nanotechnology’’ became a key word of public
interest, since even politicians and economists realized the social power of nanotechnological developments. Nanotechnology is called the technology of the next
century, coming after microtechnology. Nanotechnology unfortunately also became
a catchword for people with ambitions in science fiction. Such people threaten us
with visions of horror, as when self-replicating machines will destroy mankind.
Nevertheless, nanoscience, and consequently nanotechnology, is going to initiate a
technological impact that can probably not be compared with any other technical
development up to the present time, since it will concern all aspects of human life,
ranging from novel building materials to medicine. Of course, it is sometimes difficult to distinguish between vision and reality. However, nanotechnology has
already become part of our daily lives, even if we do not recognize this. The
most revolutionary consequences can be expected if the most valuable properties of
nanoparticles, their electronic properties, are exploited. These are of a fascinating
nature and will bring about a real technological revolution. Meanwhile, we know
the most important basic facts about nanoparticles and so can predict future
applications based on their novel properties.
This book does not, of course, deal with any kind of unsubstantiated visions, in
either a positive or a negative sense. It reports strictly on the present state of scientific development, which is impressive enough for any kind of science fiction to
be quite unnecessary.
Nanoparticles: From Theory to Application. Edited by Günter Schmid
Copyright 8 2004 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
ISBN: 3-527-30507-6
2
1 General Introduction
Semiconductor and metal nanoparticles occupy the center of scientific interest
because of their unique electronic nature: they follow quantum mechanical rules
instead of the laws of classical physics which govern bulk materials. Therefore,
small molecular clusters, mainly of the organometallic type, are not considered
further. It is also no longer helpful to strictly distinguish clusters from colloids. In
general, we call compounds of interest nanoparticles, which implies that predominantly they do not exhibit classical bulk properties, while, on the other hand, differing from molecules in so far as they in some way represent pieces of matter
related to the bulk material they originate from.
Consequently the first chapter of this book is called ‘‘Quantum Dots’’, as it deals
with those very special properties of matter in the nanosize regime. The understanding of quantum-confined electrons in very small particles is the basis for the
understanding of properties which will be described in following chapters.
Although the principal synthetic routes to nanoparticles of semiconductor or
metallic nature have been well known for some decades, some novel developments
seem worth reporting here, and for reasons of completeness classical procedures
will at least briefly be reconsidered. Syntheses of semiconductor nanoparticles of
II–VI, III–V and Ib–VI types will be followed by routes to metal nanoparticles.
The science of nanoparticles during the last decade is characterized by, among
other things, the enormous efforts which have been made to organize nanoparticles in three and two dimensions, and to some extent one dimension. Threedimensional organization of atoms, ions, or molecules is a well-known natural
process, called crystallization. Three-dimensional organization of nanoparticles is,
however, not always a simple matter, since, with some exceptions, particles of a
few up to some dozens of nanometers in size usually do not possess exactly
the same number of atoms or, consequently, the same shape. Classically, only
identical species were believed to form crystals. Then we learned that particles of
similar but not identical size and shape may be able to organize three-, even twodimensionally. These aspects become important, as we shall see, in applications of
nanoparticles in future nanoelectronic or magnetic devices, for instance in relation
to storage problems. This is why the organization of semiconductor and metal
nanoparticles will be discussed in detail.
The chapter on the various properties of nanoparticles is of central importance
in this book. With the parallel development of analytical tools, highly sophisticated
intrinsic properties of all kinds of nanoparticles became visible. Optical, electronic,
and magnetic properties of nanoparticles open up a novel world with immense
implications for future developments.
As already mentioned, biomaterials became part of nanoscience after we learned
to use biomolecules as tools to combine inorganic nanoparticles. DNA and proteins developed into valuable materials of manifold excellent utility. The combination of inorganic and biochemical building blocks led to a wide variety of novel
hybrid systems with unexpected properties. This field of nanoscience is still at the
very beginning and is expected to develop very fast in the near future.
Just like its predecessor, ‘‘Clusters and Colloids’’, this book will need a successor
in a couple of years, because of the extremely rapid development of nanoscience.
1 General Introduction
Therefore, it should be stated that this book simply gives an insight into our current, short-lived knowledge of important aspects of nanoscience.
This brings us to those parts of nanoscience that will not be considered here,
because they would overstep the limits of this book too much. Thus, we will not
consider the rapidly growing field of nanorods and nanowires. Also, nanosized
biological and supramolecular aspects will not be discussed, except in some special
cases relating to quantum dots. Nanostructured surfaces will also not be explicitly
treated, except those resulting from decoration with nanoparticles. Indeed, these
are considerable restrictions. However, it is not the purpose of this work to serve as
an encyclopedia. Rather it is an attempt to describe one important part of nanoscience that may play a decisive role in future nanotechnology.
3
4
2
Quantum Dots
Wolfgang Johann Parak, Liberato Manna, Friedrich Christian
Simmel, Daniele Gerion, and Paul Alivisatos
2.1
Introduction and Outline
In the last decade, new directions of modern research, broadly defined as ‘‘nanoscale science and technology’’, have emerged [1, 2]. These new trends involve the
ability to fabricate, characterize, and manipulate artificial structures, whose features are controlled at the nanometer level. They embrace areas of research as
diverse as engineering, physics, chemistry, materials science, and molecular biology. Research in this direction has been triggered by the recent availability of
revolutionary instruments and approaches that allow the investigation of material
properties with a resolution close to the atomic level. Strongly connected to such
technological advances are the pioneering studies that have revealed new physical
properties of matter at a level intermediate between atomic/molecular and bulk.
Materials science and technology is a field that is evolving at a very fast pace and
is currently giving the most significant contributions to nanoscale research. It is
driven by the desire to fabricate materials with novel or improved properties. Such
properties can be, for instance, strength, electrical and thermal conductivity, optical
response, elasticity, or wear resistance. Research is also evolving toward materials
that are designed to perform more complex and efficient tasks. Examples include
materials that bring about a higher rate of decomposition of pollutants, a selective
and sensitive response toward a given biomolecule, an improved conversion of
light into current, or more efficient energy storage. For such and more complex
tasks to be realized, novel materials have to be based on several components whose
spatial organization is engineered at the molecular level. This class of materials
can be defined as ‘‘nano-composites’’. They are made of assembled nanosized objects or molecules. Their macroscopic behavior arises from the combination of the
novel properties of the individual building blocks and their mutual interaction.
In electronics, the design and the assembly of functional materials and devices
based on nanoscale building blocks can be seen as the natural, inevitable evolution
of the trend toward miniaturization. The microelectronics industry, for instance, is
fabricating integrated circuits and storage media whose basic units are approachNanoparticles: From Theory to Application. Edited by Günter Schmid
Copyright 8 2004 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
ISBN: 3-527-30507-6
2.1 Introduction and Outline
ing the size of few tens of nanometers. For computers, ‘‘smaller’’ means higher
computational power at lower cost and with higher portability. However, this race
toward higher performance is driving current silicon-based electronics to the limits
of its capability [3–6]. The design of each new generation of smaller and faster
devices involves more sophisticated and expensive processing steps, and requires
the solution of new sets of problems, such as heat dissipation and device failure.
If the trend toward further miniaturization persists, silicon technology will soon
reach limits at which these problems become insurmountable. In addition to this,
scientists have found that device characteristics in very small components are
strongly altered by quantum mechanical effects. In many cases, these effects will
undermine the classical principles on which most of today’s electronic components are based. For these reasons, alternative materials and approaches are currently being explored for novel electronic components in which the laws of quantum mechanics regulate their functioning in a predictable way. Perhaps in the near
future a new generation of computers will rely on fundamental processing units
that are made only of a few atoms.
Fortunately, the advent of new methods for the controlled production of nanoscale materials has provided new tools that can be adapted for this purpose. New
terms such as nanotubes, nanowires, and quantum dots are now common jargon
of scientific publications. These objects are among the smallest man-made units
that display physical and chemical properties which make them promising candidates as fundamental building blocks for novel transistors. The advantages envisaged here are higher device versatility, faster switching speed, lower power dissipation, and the possibility of packing many more transistors on a single chip.
Prototypes of these new single nano-transistors are nowadays fabricated and
studied in research laboratories and are far from commercialization. How millions
of such components could be arranged and interconnected in complex architectures and at low cost still remains a formidable problem to be solved.
With a completely different objective, the pharmaceutical and biomedical industries try to synthesize large supramolecular assemblies and artificial devices that
mimic the complex mechanisms of nature or that can be potentially used for more
efficient diagnoses and better cures for diseases. Examples in this direction are
nanocapsules such as liposomes, embodying drugs that can be selectively released
in living organs, or bioconjugate assemblies of biomolecules and magnetic (or
fluorescent) nanoparticles that may provide faster and more selective analysis of
biotissues. These prototype systems may one day evolve into more complex nanomachines with highly sophisticated functional features able to carry out complicated tasks at the cellular level in a living body.
This chapter is not meant as a survey of the present state and future developments of nanoscale science and technology, and the list of examples mentioned
above is far from being complete. Nanoscience and nanotechnology will definitely
have a strong impact on our lives in many disparate areas. We can mention, as the
most significant examples, information technology and the telecommunications
industry, materials science and engineering, medicine. In this introductory chapter, we want to stress the point that any development in nanoscience necessarily
5
6
2 Quantum Dots
requires an understanding of the physical laws that govern matter at the nanoscale
and of how the interplay of the various physical properties of a nanoscopic system
translates into some novel behavior or into a new physical property. In this sense,
the chapter will serve as an overview of basic physical rules governing nanoscale
materials, with a particular emphasis on quantum dots, including their various
physical realizations and their possible applications. Quantum dots are the ultimate example of a solid in which all dimensions shrink down to a few nanometers.
Moreover, semiconductor quantum dots are probably the most studied nanoscale
systems.
The outline of this chapter is as follows. In Section 2.2 we try to explain with a
few examples why the behavior of nanoscale materials can be remarkably different
from that of bulk materials and from their atomic counterparts, and how quantum
mechanics can help us in rationalizing this. Following this discussion, we give a
definition of a ‘‘quantum dot’’. In Section 2.3 we follow a bottom-up approach and
give the simplified picture of a solid as being a very large molecule, where the energy levels of the individual atomic components have merged into bands. The
electronic structure of a quantum dot, being intermediate between that of the two
extreme cases of single atoms and bulk material, will then be an easier concept to
grasp. In Section 2.4 we use the model of a free-electron gas and the concept of
quantum confinement to explain what happens to a solid when its dimensions
shrink one by one. This leads us to a more accurate definition of quantum well,
quantum wire, and quantum dot. In Section 2.5 we examine in more detail the
electronic structure of quantum dots, although we try to keep the discussion at a
simple level. Section 2.6 is a brief overview of the most popular methods used to
fabricate quantum dots. Different methods lead to different varieties of quantum
dots, which can be suited for specific applications. In Section 2.7 we discuss the
optical properties of quantum dots. As they are quite unique for this class of
materials, the optical properties are probably the most important reason why the
research on quantum dots has exploded in the last decade. The discussion here
will be focused more on colloidal nanocrystal quantum dots. Electrical and transport properties are nonetheless extremely relevant, as is described in Section 2.8,
since, for instance, the addition or subtraction of a charge from a quantum dot
leads to dramatic modification of its electronic structure and of the way the dot will
handle a further addition or subtraction of a charge. This can be of fundamental
importance for future applications in electronics.
2.2
Nanoscale Materials and Quantum Mechanics
2.2.1
Nanoscale Materials as Intermediate between Atomic and Bulk Matter
Nanoscale materials frequently show behavior which is intermediate between that
of a macroscopic solid and that of an atomic or molecular system. Consider, for
2.2 Nanoscale Materials and Quantum Mechanics
instance, the case of an inorganic crystal composed of few atoms. Its properties
will be different from those of a single atom, but we cannot imagine that they will
be the same as those of a bulk solid. The number of atoms on the crystal’s surface,
for instance, is a significant fraction of the total number of atoms, and therefore
will have a large influence on the overall properties of the crystal. We can easily
imagine that this crystal might have a higher chemical reactivity than the corresponding bulk solid and that it will probably melt at lower temperatures. Consider
now the example of a carbon nanotube, which can be thought of as a sheet of
graphite wrapped in such a way that the carbon atoms on one edge of the sheet are
covalently bound to the atoms on the opposite edge of the sheet. Unlike its individual components, a carbon nanotube is chemically extremely stable because the
valences of all its carbon atoms are saturated. Moreover, we would guess that
carbon nanotubes can be good conductors because electrons can freely move along
these tiny, wire-like structures. Once again, we see that such nanoscopic objects
can have properties which do not belong to the realm of their larger (bulk) or
smaller (atoms) counterparts. However, there are many additional properties specific to such systems which cannot easily be grasped by simple reasoning. These
properties are related to the sometimes counterintuitive behavior that charge carriers (electrons and holes) can exhibit when they are forced to dwell in such structures. These properties can only be explained by the laws of quantum mechanics.
2.2.2
Quantum Mechanics
A fundamental aspect of quantum mechanics is the particle-wave duality, introduced by De Broglie, according to which any particle can be associated with a
matter wave whose wavelength is inversely proportional to the particle’s linear
momentum. Whenever the size of a physical system becomes comparable to the
wavelength of the particles that interact with such a system, the behavior of the
particles is best described by the rules of quantum mechanics [7]. All the information we need about the particle is obtained by solving its Schrödinger equation.
The solutions of this equation represent the possible physical states in which the
system can be found. Fortunately, quantum mechanics is not required to describe
the movement of objects in the macroscopic world. The wavelength associated with
a macroscopic object is in fact much smaller than the object’s size, and therefore
the trajectory of such an object can be excellently derived using the principles of
classical mechanics. Things change, for instance, in the case of electrons orbiting
around a nucleus, since their associated wavelength is of the same order of magnitude as the electron-nucleus distance.
We can use the concept of particle-wave duality to give a simple explanation of
the behavior of carriers in a semiconductor nanocrystal. In a bulk inorganic semiconductor, conduction band electrons (and valence band holes) are free to move
throughout the crystal, and their motion can be described satisfactorily by a linear
combination of plane waves whose wavelength is generally of the order of nanometers. This means that, whenever the size of a semiconductor solid becomes
7
8
2 Quantum Dots
comparable to these wavelengths, a free carrier confined in this structure will
behave as a particle in a potential box [8]. The solutions of the Schrödinger equation in such case are standing waves confined in the potential well, and the energies associated with two distinct wavefunctions are, in general, different and
discontinuous. This means that the particle energies cannot take on any arbitrary
value, and the system exhibits a discrete energy level spectrum. Transitions between any two levels are seen as discrete peaks in the optical spectra, for instance.
The system is then also referred to as ‘‘quantum confined’’. If all the dimensions of
a semiconductor crystal shrink down to a few nanometers, the resulting system is
called a ‘‘quantum dot’’ and will be the subject of our discussion throughout this
chapter. The main point here is that in order to rationalize (or predict) the physical
properties of nanoscale materials, such as their electrical and thermal conductivity
or their absorption and emission spectra, we need first to determine their energy
level structure.
For quantum-confined systems such as quantum dots, the calculation of the
energy structure is traditionally carried out using two alternative approaches. One
approach has just been outlined above. We take a bulk solid and we study the evolution of its band structure as its dimensions shrink down to a few nanometers.
This method will be described in more detail later (Section 2.4). Alternatively, we
can start from the individual electronic states of single isolated atoms as shown in
Section 2.3 and then study how the energy levels evolve as atoms come closer and
start interacting with each other.
2.3
From Atoms to Molecules and Quantum Dots
From the point of view of a chemist, the basic building blocks of matter are atomic
nuclei and electrons. In an atom, electrons orbit around the single nucleus, and the
number of electrons depends on the element. In the simplest case of the hydrogen
atom, one electron orbits around one proton. The electronic states of the hydrogen
atom can be calculated analytically [9, 10]. As soon as more than one electron
is involved, however, the calculation of the energy levels becomes more complicated, since, in addition to the interaction between the nucleus and the electron, now
also electron-electron interactions have to be taken into account. Although the energy states of many-electron atoms can no longer be derived analytically, approximations such as the Hartree-Fock method exist [10]. Each electron can be ascribed
to an individual orbit, called an atomic orbital (AO), with an associated discrete
energy level. Depending on the angular moment of the orbit, AOs can have
spherical (s-orbital), club-like ( p-orbital) or a more complicated (d-, f -orbitals)
shape. The eight valence electrons of a neon atom, for example, occupy one s- and
three p-orbitals around the nucleus, one spin up and one spin down per orbit [10],
where the energy level of the s-orbital is lower than that of the p-orbitals. In accordance with the rules of quantum mechanics, the energy levels are discrete.
The next bigger structure, obtained from the combination of several atoms, is
2.3 From Atoms to Molecules and Quantum Dots
energy
σ*
p
conduction
band
sp3
energy gap
s
atom
σ
molecule
valence
band
quantum
dot
number of connected atoms
Electronic energy levels depending on
the number of bound atoms. By binding more
and more atoms together, the discrete energy
levels of the atomic orbitals merge into energy
bands (here shown for a semiconducting
Fig. 2-1
bulk solidstate body
material) [16]. Therefore semiconducting
nanocrystals (quantum dots) can be regarded
as a hybrid between small molecules and bulk
material.
the molecule. Now electrons orbit collectively around more than one nucleus. In a
molecule, electrons that are responsible for the covalent bonds between individual
atoms can no longer be ascribed to one individual atom, but they are ‘‘shared’’. In
methane (CH4 ), for instance, each of the four sp 3 atomic orbitals of the central
carbon atom is linearly combined with the s orbital of a hydrogen atom to form a
bonding (s) and an anti-bonding (s ) orbital, respectively [9]. Since these orbitals
are ‘‘shared’’ between the atoms, they are called molecular orbitals (MO, see Figure
2-1). Only the lowest energy (bonding) orbitals are occupied, and this explains the
relative stability of methane [10]. Using the same principle, it is possible to derive
the electronic structure of more complex systems such as large molecules or
atomic clusters. When combining atoms to form a molecule, we start from discrete
energy levels of the atomic orbitals and we still end up obtaining discrete levels for
the molecular orbitals [9].
When the size of a polyatomic system becomes progressively larger, the calculation of its electronic structure in terms of combinations of atomic orbitals becomes
unfeasible [11–13]. However, simplifications arise if the system under study is a
periodic, infinite crystal. The electronic structure of crystalline solids can be in fact
described in terms of periodic combinations of atomic orbitals (Bloch functions)
[14, 15]. In this model, perfect translational symmetry of the crystal structure is
assumed, and contributions from the surface of the crystal are neglected by assuming an infinite solid (periodic boundary conditions). Electrons are described as
a superposition of plane waves extended throughout the solid. As opposed to the
case of atoms and molecules, the energy structure of a solid no longer consists of
discrete energy levels, but rather of broad energy bands [14, 15], as sketched in
Figure 2-1. Every band can be filled only with a limited amount of charge carriers.
9
10
2 Quantum Dots
In very small crystals of nanometer dimensions, so called nanocrystals, the assumptions of translational symmetry and infinite size of the crystal are no longer
valid, and thus these systems cannot be described with the same model used for a
bulk solid. We can imagine indeed that the electronic structure of a nanocrystal
should be something intermediate between the discrete levels of an atomic system
and the band structure of a bulk solid. This can be evidenced from Figure 2-1: the
energy levels of a nanocrystal are discrete, their density is much larger, and their
spacing is smaller than for the corresponding levels of one atom or a small atomic
cluster. Because of their discrete energy levels, such structures are called also
quantum dots. The concept of energy bands and band gap can still be used. Highest occupied atomic levels of the atomic (or ionic) species interact with each other
to form the valence band of the nanocrystal. Similarly, lowest unoccupied levels
combine to form the conduction band of the nanocrystal. The energy gap between
the valence and conduction bands results in the band gap of the nanocrystal. As an
example, consider a metallic quantum dot. Its level spacing at the Fermi level is
roughly proportional to @EF =N, where N is the number of electrons in the quantum dot. Given that EF is a few eV and that N is close to 1 per atom, the band gap
of a metallic quantum dot becomes observable only at very low temperatures.
Conversely, in the case of semiconductor quantum dots, the band gap is larger and
its effects can be observed at room temperature. The size-tunable fluorescence
emission of CdSe quantum dots in the visible region of the spectrum is for instance a very explanatory illustration of the presence of a size-dependent band gap.
2.4
Shrinking Bulk Material to a Quantum Dot
In this chapter we will use the concept of quantum confinement of carriers in a
solid to derive a more detailed description of the electronic band structure in a lowdimensional solid. This description, although more elaborate than the one just
given above, is indeed more powerful and will underline the general physics of a
solid when its dimensions shrink one by one down to few nanometers. We will
start first with an elementary model of the behavior of electrons in a bulk solid.
This model will then be adapted to the case of confined carriers.
2.4.1
Three-Dimensional Systems (Bulk Material)
We now consider the case of a three-dimensional solid of size dx ; d y ; d z , containing
N free electrons. ‘‘Free’’ means that these electrons are delocalized and thus not
bound to individual atoms. Furthermore, we will make the assumption that the
interactions between the electrons, as well as the interactions between the electrons and the crystal potential, can be neglected as a first approximation. Such a
model system is called ‘‘free-electron gas’’ [14, 15].
2.4 Shrinking Bulk Material to a Quantum Dot
Astonishingly, this oversimplified model still captures many of the physical aspects of real systems. From more complicated theories it has been learnt that many
of the expressions and conclusions from the free-electron model remain valid as a
first approximation even when one takes electron-crystal and electron-electron interactions into account. In many cases it is sufficient to replace the free-electron
mass m by an ‘‘effective’’ mass m , which implicitly contains the corrections for
the interactions. To keep the story simple, we proceed with the free-electron picture. In the free-electron model, each electron in the solid moves with a velocity
~
v ¼ ðvx ; vy ; vz Þ. The energy of an individual electron is then just its kinetic energy:
1 2 1
v ¼ mðvx 2 þ vy 2 þ vz 2 Þ
E ¼ m~
2
2
ð1Þ
According to Pauli’s exclusion principle, each electron must be in a unique quantum state. Since electrons can have two spin orientations (m s ¼ þ 12 and m s ¼ 12),
only two electrons with opposite spins can have the same velocity ~
v. This case is
analogous to the Bohr model of atoms, in which each orbital can be occupied by
two electrons at maximum. In solid-state physics, the wavevector ~
k ¼ ðkx ; k y ; kz Þ of
a particle is more frequently used instead of its velocity to describe the particle’s
state. Its absolute value k ¼ j~
kj is the wavenumber. The wavevector ~
k is directly
proportional to the linear momentum ~
p and thus also to the velocity ~
v of the
electron:
~
p ¼ m~
v¼
h~
k
2p
ð2Þ
The scaling constant is the Planck constant h, and the wavenumber is related to
the wavelength l associated with the electron through the De Broglie relation [14,
15]:
Gk ¼ j~
kj ¼ G
2p
l
ð3Þ
The wavelengths l associated with the electrons traveling in a solid are typically of
the order of nanometers1), much smaller than the dimensions of an ordinary solid.
The calculation of the energy states for a bulk crystal is based on the assumption
of periodic boundary conditions. Periodic boundary conditions are a mathematical
trick to ‘‘simulate’’ an infinite ðd ! yÞ solid. This assumption implies that the
conditions at opposite borders of the solid are identical. In this way, an electron
that is close to the border does not really ‘‘feel’’ the border. In other words, the
electrons at the borders ‘‘behave’’ exactly as if they were in the bulk. This con1) In fact, the wavelength depends on the
electron density. The wavelength for electrons
in metals is typically around 10 nm; in
semiconductors it may vary between 10 nm
and 1 mm.
11
12
2 Quantum Dots
V(x)
λ = dx/3
k = ±6 π /dx
n=3
λ = dx/2
k = ±4 π /dx
ψ(x)
n=2
λ = dx
k = ±2 π /dx
n=1
n=0
-½ dx
0
½ dx
λ=∞
k=0
x
Periodic boundary conditions (only
drawn for the x-dimension) for a free-electron
gas in a solid with thickness d. The idea of
periodic boundary conditions is to mathematically ‘‘simulate’’ an infinite solid. Infinite
extension is similar to an object without any
borders. This means that a particle close to the
‘‘border’’ must not be affected by the border,
but ‘‘behaves’’ exactly as it were in the bulk.
This can be realized by using a wavefunction
cðxÞ that is periodic within the thickness d of
the solid. Any electron that leaves the solid
from its right boundary would reenter under
Fig. 2-2
exactly the same conditions on its left side. For
the electron the borders are quasi-nonexistent.
The probability density jcðxÞj 2 is the probability
that an electron is at the position x in the
solid. Different states for the electrons
ðn ¼ 0; 1; 2; . . .Þ have different wavefunctions. l
is the De Broglie wavelength of the electrons
and k is their corresponding wavenumber. A
‘‘real’’ bulk solid can be approximated by an
infinite solid ðd ! yÞ and its electronic states
in k-space are quasi-continuously distributed:
Dk ¼ 2p/dx ! 0.
dition can be realized mathematically by imposing the following condition to
the electron wavefunctions: cðx; y; zÞ ¼ cðx þ dx ; y; zÞ, cðx; y; zÞ ¼ cðx; y þ d y ; zÞ,
and cðx; y; zÞ ¼ cðx; y; z þ d z Þ. In other words, the wavefunctions must be periodic with a period equal to the whole extension of the solid [15, 17]. The solution of
the stationary Schrödinger equation under such boundary conditions can be factorized into the product of three independent functions cðx; y; zÞ ¼ cðxÞcð yÞcðzÞ ¼
A expðikx xÞ expðik y yÞ expðikz zÞ. Each function describes a free electron moving
along one Cartesian coordinate. In the argument of the functions, kx; y; z is equal to
GnDk ¼ Gn2p=dx; y; z and n is an integer [14, 15, 17]. These solutions are waves
that propagate along the negative and the positive direction, for kx; y; z > 0 and
kx; y; z < 0, respectively. An important consequence of the periodic boundary conditions is that all the possible electronic states in the ~
k space are equally distributed. There is an easy way of visualizing this distribution
case of a
in the ideal
one-dimensional free-electron gas: there are two electrons m s ¼ G 12 in the state
kx ¼ 0 ðvx ¼ 0Þ, two electrons in the state kx ¼ þDk ðvx ¼ þDvÞ, two electrons in
the state kx ¼ Dk ðvx ¼ DvÞ, two electrons in the state kx ¼ þ2Dk ðvx ¼ þ2DvÞ,
and so on.
For a three-dimensional
bulk material we can follow an analogous scheme. Two
electrons m s ¼ G 12 can occupy each of the states ðkx ; k y ; kz Þ ¼ ðGnx Dk;Gn y Dk;
2.4 Shrinking Bulk Material to a Quantum Dot
Gnz DkÞ, again with nx; y; z being an integer. A sketch of this distribution is shown
in Figure 2-3. We can easily visualize the occupied states in ~
k-space because all
these states are included into a sphere whose radius is the wavenumber associated
with the highest energy electrons. At the ground state, at 0 K, the radius of the
sphere is the Fermi wavenumber kF (Fermi velocity vF ). The Fermi energy
EF z kF 2 is the energy of the last occupied electronic state. All electronic states
with an energy E a EF are occupied, whereas all electronic states with higher
energy E > EF are empty. In a solid, the allowed wave numbers are separated by
Dk ¼ Gn2p=dx; y; z . In a bulk material dx; y; z is large, and so Dk is very small. Then
the sphere of states is filled quasi-continuously [15].
We need now to introduce the useful concept of the density of states D3d ðkÞ,
which is the number of states per unit interval of wavenumbers. From this definition, D3d ðkÞDk is the number of electrons in the solid with a wavenumber between
k and k þ Dk. If we know the density of states in a solid we can calculate, for instance, the total number of electrons having wavenumbers less
Ð k than a given k max ,
which we will call Nðk max Þ. Obviously, Nðk max Þ is equal to 0 max D3d ðkÞ dk. In the
ground state of the solid, all electrons have wavenumbers k a kF , where kF is the
Fermi wavenumber. Since in a bulk solid the states are homogeneously distributed
in ~
k-space, we know that the number of states between k and k þ Dk is proportional
to k 2 Dk (Figure 2-3). This can be visualized in the following way. The volume in
three-dimensional ~
k-space varies with k 3 . If we only want to count the number of
states with a wavenumber between k and k þ Dk, we need to determine the volume
of a spherical shell with radius k and thickness Dk. This volume is proportional to
product of the surface of the sphere (which varies as k 2 ) with the thickness of the
shell (which is Dk). D3d ðkÞDk is thus proportional to k 2 Dk, and in the limit when Dk
approaches zero, we can write:
D3d ðkÞ ¼
dNðkÞ
z k2
dk
ð4Þ
Instead of knowing the density of states in a given interval of wavenumbers it is
more useful to know the number of electrons that have energies between E and
2
E þp
DE.
ffiffiffi From Eqs. (1) and (2) we know
pffiffiffi that EðkÞ is proportional to k , and thus
k z E . Consequently, dk=dE z 1= E . By using Eq. (4), we obtain for the density
of states for a three-dimensional electron gas [17]:
D3d ðEÞ ¼
pffiffiffi pffiffiffi
dNðEÞ dNðkÞ dk
¼
z E 1= E z E
dE
dk dE
ð5Þ
This can be seen schematically in Figure 2-3. With Eq. (5) we conclude our simple
description of a bulk solid. The possible states in which an electron can be found
are quasi-continuous. The density of states varies with the square root of the energy. More details about the free-electron gas model and more refined descriptions
of electrons in solids can be found in any solid state physics textbook [14].
13
14
2 Quantum Dots
b)
a)
(kx, ky, kz)
Dky
kz
kz
z
x
c)
kx
dx
y
kx
E(kx,y,z)
Dkx,y,z → 0
d)
kx,y,z
Electrons in a three-dimensional bulk
solid [15]. (a) Such a solid can be modeled
as an infinite crystal along all three dimensions x; y; z. (b) The assumption of periodic
boundary conditions yields standing waves
as solutions for the Schr€
odinger equation for
free electrons. The associated wavenumbers
ðkx ; k y ; kz Þ are periodically distributed in the
reciprocal k-space [17]. Each of the dots shown
in the figure represents a possible electronic
state ðkx ; k y ; kz Þ. Each state in k-space can be
only occupied by two electrons. In a large solid
the spacing Dkx; y; z between individual electron
states is very small, and therefore the k-space
is quasi-continuously filled with states. A
sphere with radius kF includes all states with
k ¼ ðkx 2 þ k y 2 þ kz 2 Þ 1/2 < kF . In the ground
Fig. 2-3
ky
ky
D3d (E)
E
state, at 0 K, all states with k < kF are occupied by two electrons, and the other states
are empty. Since the k-space is homogeneously filled with states, the number of states
within a certain volume varies with k 3 . (c) Dispersion relation for free electrons in a threedimensional solid. The energy of free electrons
varies with the square of the wavenumber,
and its dependence on k is described by a
parabola. For a bulk solid the allowed states
are quasi-continuously distributed and the
distance between two adjacent states (here
shown as points) in k-space is very small.
(d) Density of states D3d for free electrons in
a three-dimensional system. The allowed
energies are quasi-continuous and their density
varies with the square root of the energy E 1/2.
2.4.2
Two-Dimensional Systems
We now consider a solid that is fully extended along the x- and y-directions, but
whose thickness along the z-direction (d z ) is only a few nm (see Figure 2-5). Free
electrons can still move freely in the x-y plane. However, movement in the zdirection is now restricted. Such a system is called a 2-dimensional electron gas
(2DEG) [18]. As mentioned in Section 2.2, when one or more dimensions of a
solid become smaller than the De Broglie wavelength associated with the free
charge carriers, an additional contribution of energy is required to confine the
2.4 Shrinking Bulk Material to a Quantum Dot
component of the motion of the carriers along this dimension. In addition, the
movement of electrons along such a direction becomes quantized. This situation is
shown in Figure 2-4. No electron can leave the solid, and electrons that move in
the z-direction are trapped in a ‘‘box’’. Mathematically this is described by infinitely high potential wells at the border z ¼ G 12 d z.
The solutions for the particle-in-a-box situation can be obtained by solving the
one-dimensional Schrödinger equation for an electron in a potential VðzÞ, which
is zero within the box but infinite at the borders. As can be seen in Figure 2-4, the
solutions are stationary waves with energies2) Enz ¼ p 2 kz 2 =2m ¼ h 2 kz 2 =8p 2 m ¼
h 2 nz 2 =8md z 2 , nz ¼ 1; 2; . . . [9, 17]. This is similar to states kz ¼ nz Dkz with Dkz ¼
p=d z . Again, each of these states can be occupied at maximum by two electrons.
Let us compare the states in the k-space for three- and two dimensional materials
(Figures 2-3 and 2-5). For a two-dimensional solid that is extended in the x-y-plane
only discrete values are allowed for kz. The thinner the solid in the z-direction, the
larger is the spacing Dkz between these allowed states. On the other hand, the distribution of states in the kx -k y plane remains quasi-continuous. Therefore one can
describe the possible states in the k-space as planes parallel to the kx - and k y -axes,
with a separation Dkz between the planes in the kz -direction. We can number the
2) The particle-in-a-box approach (Figure 2-4)
looks similar to the case of the periodic
boundary conditions (Figure 2-2). There are
indeed important differences between the
two cases. Periodic boundary conditions
‘‘emulate’’ an infinite solid. A quantum
mechanical treatment of this problem yields
propagating waves that are periodic within
the solid. Such waves can be seen as the
superposition of plane waves. For an idealized
one-dimensional solid, with boundaries fixed
at x ¼ Gd=2, a combination of plane waves
can be, for instance, cðxÞ ¼ A expðikxÞ þ
B expðikxÞ with k ¼ n2p=d. Written in
another way, the solutions are of the type
expðikxÞ, with k ¼ Gn2p=d. The solutions for
k ¼ þn2p=d and k ¼ n2p=d are linearly
independent. The waves expðþin2px=dÞ propagate to the right, the waves expðin2px=dÞ
to the left side of the solid. Neither wave feels
the boundaries. Since expðikxÞ ¼ cosðkxÞ þ
i sinðkxÞ and expðikxÞ ¼ cosðkxÞ i sinðkxÞ,
we also can write cðxÞ ¼ C sinðkxÞ þ
D cosðkxÞ with k ¼ n2p=d as solutions. The
only constraint here is that the wavefunction
must be periodic throughout the solid. The
state with wavenumber k ¼ 0 is a solution,
since C sinð0Þ þ D cosð0Þ ¼ D 0 0. Therefore the state with the lowest kinetic energy is
E z k 2 ¼ 0 for k ¼ 0. The individual states in
k-space are very close to each other because
Dk ¼ 2p=d tends to 0 when d increases. On
the other hand, the particle-in-a-box model
describes the case in which the motion of the
electrons is confined along one or more directions. Outside the box the probability of finding an electron is zero. For a one-dimensional
problem the solutions are standing waves of
the type cðxÞ ¼ A sinðkxÞ with k ¼ np=d.
There is only one solution of this type. The
function cðxÞ ¼ B sinðkxÞ can be written
as cðxÞ ¼ B sinðkxÞ and therefore is still of
the type cðxÞ ¼ A sinðkxÞ. Because of the
boundary conditions cðx ¼ Gd=2Þ ¼ 0 there
is no solution of the type cðxÞ ¼ B cosðkxÞ.
Since the standing wave is confined to the
box, there is only the solution k ¼ þnp=d > 0.
For a small box the energy states are far apart
from each other in k-space, and the distribution of states and energies is discrete. An
important difference with respect to the
extended solid is the occurrence of a finite
zero-point energy [9]. There is no solution for
k ¼ 0, since cð0Þ ¼ A sinð0Þ ¼ 0. Therefore
the energy of the lowest possible state ðn ¼ 1Þ
is equal to E ¼ h 2 =8md 2 , i.e. k ¼ p=d. This
energy is called zero-point energy and is a
purely quantum mechanical effect. It can be
understood as the energy that is required to
‘‘confine’’ the electron inside the box. For a
large box the zero-point energy tends to zero.
However, for small boxes this energy becomes
significant as it varies with the square of the
reciprocal of the box size d 2 .
15
16
2 Quantum Dots
∞
V(z)
∞
λ = 2/3dz
k = 3 π /dz
n=3
ψ(z)
λ = dz
k = 2 π /dz
n=2
n=1
0
-½ dz
λ = 2dz
k = 1 π /dz
x
0
½ dz
Particle-in-a-box model for a free
electron moving along in the z-axis. The
movement of electrons in the z-direction is
limited to a ‘‘box’’ with thickness d: since
electrons cannot ‘‘leave’’ the solid (the box),
their potential energy VðxÞ is zero within the
Fig. 2-4
solid, but is infinite at its borders. The
probability density jcðzÞj 2 is the probability
that an electron is located at position x in the
solid. Different states for the electrons
ðn ¼ 1; 2; . . .Þ differ in their wavefunction.
individual planes with nz . Since within one plane the number of states is quasicontinuous, the number of states is proportional to the area of the plane. This
means that the number of states is proportional to k 2 ¼ kx 2 þ k y 2 . The number of
states in a ring with radius k and thickness Dk is therefore proportional to k Dk.
Integration over all rings yields the total area of the plane in k-space. Here, in
contrast to the case of a three-dimensional solid, the density of states varies linearly
with k:
D2d ðkÞ ¼
dNðkÞ
zk
dk
ð6Þ
In the ground state, all states with k a kF are occupied by two electrons. We now
want to know how many states exist for electrons that have energies between E and
2
E þ DE. FrompEqs.
know
ffiffiffi (1) and (2) we p
ffiffiffi the relation between k and E: EðkÞ z k
and thus k z E and dk=dE z 1= E . By using Eq. (6) we obtain the density of
states for a 2-dimensional electron gas, see also Figure 2-5 [17].
D2d ðEÞ ¼
pffiffiffi
pffiffiffi
dNðEÞ dNðkÞ dk
¼
z E 1= E z 1
dE
dk dE
ð7Þ
The density of electronic states in a two-dimensional solid is therefore remarkably
different from the three-dimensional case. The spacing between the allowed energy levels in the bands increases, because fewer levels are now present. As soon as
one dimension is reduced to nanometer size, dramatic changes due to quantum
confinement occur, as, for example, the non-negligible zero-point energy. In two-
2.4 Shrinking Bulk Material to a Quantum Dot
b)
a)
(kx, ky)
Dky
kF
z
kz
dx
kz
y
x
kx
c)
kx
ky
d)
E(k z)
E(k x,y)
ky
D2d (E)
nz=4
nz=4
nz=3
nz=3
nz=2
nz=2
k x,y
Dk x,y → 0
nz=1
Electrons in a two-dimensional
system. (a) A two-dimensional solid is
(almost) infinitely extended in two dimensions
(here x; y), but is very thin along the third
dimension (here denoted as z), which is
comparable to the De Broglie wavelength of a
free electron (d z ! l.). (b) Electrons can still
move freely along the x- and y-directions. The
wavefunctions along such directions can be
found again by assuming periodic boundary
conditions. kx and k y states are quasicontinuously distributed in k-space. The
movement of electrons in the z-direction is
restricted and electrons are confined to a
‘‘box’’. Only certain quantized states are
allowed along this direction. For a discrete
kz -state, the distribution of states in threedimensional k-space can be described as a
series of planes parallel to the kx - and k y -axes.
For each discrete kz -state, there is a separate
plane parallel to the kx and to the k y -axes.
Here only one of those planes is shown. The
kx - and k y -states within one plane are quasicontinuous, since Dkx; y ¼ 2p/dx; y ! 0. The
distance between two planes for two separate kz -states is large, since Dkz ¼ p/d z g 0.
For each kz -value the kx - and k y states are
homogeneously distributed on the kx -k y -plane
Fig. 2-5
nz=1
kz
E
Dk z
[17]. The number of states within this plane is
therefore proportional to the area of a disk
around kx ¼ k y ¼ 0. This means that the
number of states for a certain wavenumber
varies with k 2 . In the ground state all states
with k a kF are occupied with two electrons,
while the remaining states are empty. (c) Free
electrons have a parabolic dispersion relation
ðEðkÞ z k 2 Þ. The energy levels Eðkx Þ and Eðk y Þ
for the electron motion along the x- and the
y-directions are quasi-continuous (they are
shown here as circles). The wavefunction cðzÞ
at the border of a small ‘‘box’’ must be zero,
leading to standing waves inside the box.
This constraint causes discrete energy levels
Eðkz Þ for the motion along the z-direction.
Electrons can only occupy such discrete states
(nz1 ; nz2 ; . . . ; shown here as circles). The
position of the energy levels now changes with
the thickness of the solid in the z-direction,
or in other words with the size of the ‘‘box’’.
(d) Density of states for a two-dimensional
electron gas. If electrons are confined in one
direction ðzÞ but can move freely in the other
two directions ðx; yÞ, the density of states for a
given kz -state ðnz ¼ 1; 2; . . . :Þ does not depend
on the energy E.
17
18
2 Quantum Dots
dimensional materials the energy spectrum is still quasi-continuous, but the density of states now is a step function [17, 19].
The quantum-mechanical behavior of electrons in a two-dimensional solid is the
origin of many important physical effects. With recent progress in nanoscience
and technology, the fabrication of two-dimensional structures has become routine.
2D systems are usually formed at interfaces between different materials or in layered systems in which some of the layers may be only a few nanometers thick.
Structures like this can be grown, for example, by successive deposition of the
individual layers by molecular beam epitaxy. In such geometry, charge carriers
(electrons and holes) can move freely parallel to the semiconductor layer, but their
movement perpendicular to the interface is restricted. The study of these nanostructures led to the discovery of remarkable 2-dimensional quantized effects, such
as the Integer and the Fractional Quantum Hall Effect [20–23].
2.4.3
One-Dimensional Systems (Quantum Wires)
Let us now consider the case in which the solid also shrinks along a second ( y)
dimension. Now electrons can only move freely in the x-direction, and their
motion along the y- and z-axes is restricted by the borders of the solid (see Figure
2-6). Such a system is called a quantum wire or – when electrons are the charge
carriers – a one-dimensional electron system (1DES). The charge carriers and the
b)
a)
kz
z
x
kx = 2Dkx
y
kx
c)
Dk x → 0
kx
(a) One-dimensional solid. (b) The
allowed ðkx ; k y ; kz Þ-states can be visualized
as lines parallel to the kx -axes in the threedimensional k-space. In this figure only one
line is shows as an example. Within each line,
the distribution of states is quasi-continuous,
since Dkx ! 0. The arrangement of the
individual lines is discrete, since only certain
discrete k y - and kz -states are allowed. (c) This
Dkx
ky
d)
E(k y,z)
E(kx)
Fig. 2-6
kx = Dkx
kx = -Dkx
kx = 0
D1d (E)
k y,z
E
Dk y,z
can also be seen in the dispersion relations.
Along the kx -axes the energy band Eðkx ; k y ; kz Þ
is quasi-continuous, but along the k y - and kz axes only certain energies exist. (d) The density
of states within one line along the kx -axes is
proportional to E1/2 . Each of the hyperbolas
shown in the D1d -diagram corresponds to an
individual ðk y ; kz Þ-state.
2.4 Shrinking Bulk Material to a Quantum Dot
excitations now can move only in one dimension and occupy quantized states in
the other two dimensions.
The states of a one-dimensional solid now can be obtained by methods analogous to those described for the three- and two-dimensional materials. In the xdirection electrons can move freely, and again we can apply the concept of periodic
boundary conditions. This yields a quasi-continuous distribution of states parallel
to the kx -axis as well as a quasi-continuous distribution of the corresponding energy levels. Electrons are confined along the remaining directions and their states
can be derived from the Schrödinger equation for a particle-in-a-box potential.
Again, this yields discrete k y and kz -states. We can now visualize all possible states
as lines parallel to the kx -axis. The lines are separated by discrete intervals along k y
and kz , but within each line the distribution of kx states is quasi-continuous (Figure 2-6). We can count the number of states along one line by measuring the
length of the line. The number of states is therefore proportional to k ¼ kx . Hence,
the number of states with wavenumbers in the interval between k and k þ Dk is
proportional to Dk:
D1d ðkÞ ¼
dNðkÞ
z1
dk
ð8Þ
In the ground state, all states with k a kF are occupied by two electrons. From Eqs.
(1) and (2) we p
know
k and E for free electron is: EðkÞ z k 2
ffiffiffi that the relation
pbetween
ffiffiffi
and thus k z E and dk=dE z 1= E . By using Eq. (8) we obtain the density of
states for a 1-dimensional electron gas:
D1d ðEÞ ¼
pffiffiffi
pffiffiffi
dNðEÞ dNðkÞ dk
¼
z 1 1= E z 1= E
dE
dk dE
ð9Þ
The density of states is depicted in Figure 2-6. In one-dimensional systems the
density of states has a E 1=2 -dependence and thus exhibits singularities near the
band edges [17]. Each of the hyperbolas contains a continuous distribution of kx
states but only one discrete k y - and kz -state.
The quantization of states in two dimensions has important consequences for
the transport of charges. Electrons can only flow freely along the x-axes but are
limited to discrete states in the y- and z-directions. Therefore they are only transported in discrete ‘‘conductivity channels’’. This may be of considerable importance
for the microelectronics industry. If the size of electronic circuits is reduced more
and more, at one point the diameter of wires will become comparable to the De
Broglie wavelength of the electrons. The wire will then exhibit the behavior of a
quantum wire. Quantum aspects of 1D transport were first observed in so-called
quantum point contacts, which were lithographically defined in semiconductor
heterostructures [24, 25]. More recent examples of such 1D wires include short
organic semiconducting molecules [26–31], inorganic semiconductor and metallic
nanowires [32–37] or break junctions [38–40]. A particular role is played by carbon
nanotubes [27, 41–47]. Carbon nanotubes have been extensively studied both as
model systems for one-dimensional confinement and for potential applications,
such as electron emitters [48].
19
20
2 Quantum Dots
2.4.4
Zero-Dimensional Systems (Quantum Dots)
When charge carriers and excitations are confined in all three dimensions, the
system is called a ‘‘quantum dot’’. The division is somewhat arbitrary, since, for
instance, clusters made of very few atoms are not necessarily considered as
quantum dots. Although clusters are smaller than the De Broglie wavelength, their
properties depend critically on their exact number of atoms. Large clusters have a
well-defined lattice and their properties no longer depend critically on their exact
number of atoms. We will therefore use the term ‘‘quantum dots’’ to refer to such
systems [49–59].
In a quantum dot, the movement of electrons is confined in all three dimensions
and there are only discrete ðkx ; k y ; kz Þ-states in the k-space. Each individual state in
k-space can be represented by a point. The final consequence is that only discrete
energy levels are allowed, and they can be seen as delta-peaks in the distribution
D0d ðEÞ. As we can see, the energy bands converge to atom-like energy states, with
the oscillator strength compressed into a few transitions. This change is most
dramatic at the edges of the bands and so influences semiconductors more than
metals. In semiconductors the electronic properties are in fact strongly related to
the transitions between the edges of the valence band and the conduction band,
respectively. In addition to the discreteness of the energy levels, we want to stress
again the occurrence of a finite zero-point energy. In a dot, even in the ground
state, electrons have energies larger than bulk electrons at the conduction band
edge. These points will be discussed in more detail in the next section.
a)
x
c)
b)
z
y
kz
kx
(kx=0, ky=0, kz=0)
ky
d)
E(kx,y,z)
D0d(E)
kx,y,z
E
Dkx,y,z
A zero-dimensional solid. (a) The
solid is shrunk in all three dimensions to a
thickness that is comparable to the De Broglie
wavelength of its charge carriers. (b) Because
of such confinement, all states ðkx ; k y ; kz Þ are
discrete points in the three-dimensional k-
Fig. 2-7
space. (c) Only discrete energy levels are
allowed. (d) The one-dimensional density of
states D0d ðEÞ contains delta peaks, which
correspond to the individual states. Electrons
can occupy only states with these discrete
energies.
2.5 Energy Levels of a (Semiconductor) Quantum Dot
2.5
Energy Levels of a (Semiconductor) Quantum Dot
In this section we will describe in more detail a zero-dimensional solid. Since
many quantum effects are more pronounced in semiconductors compared to metals, we will focus on the case of a semiconducting material. In Section 2.4 we have
described how the properties of a free-electron gas change when the dimensions of
the solid are reduced. The model of the free-electron gas does not include the
‘‘nature’’ of the solid. However, from a macroscopic point of view we distinguish
between metals, semiconductors and insulators [14]. The model of a free-electron
gas describes relatively well the case of electrons in the conduction band of metals.
On the other hand, electrons in an insulating material are only poorly described by
the free-electron model. In order to extend the model of free electrons for semiconducting materials the concept of a new charge carrier, the hole, was introduced
[15]. If one electron from the valence band is excited to the conduction band, the
‘‘empty’’ electronic state in the valence band is called a hole. Some basic properties
of semiconducting materials can be described by the model of free electrons and
free holes. The energy bands for electrons and holes are separated by a band gap
[14, 15]. The dispersion relations for the energy of electrons and holes in a semiconductor are parabolic at first approximation. This approximation holds true only
for electrons (holes) occupying the levels that lie at the bottom (top) of the conduction (valence) band. Each parabola represents a quasi-continuous set of electron
(hole) states along a given direction in k-space. The lowest unoccupied energy band
and the highest occupied energy band are separated by an energy gap Eg (bulk), as
shown in Figure 2-8. The band gap for a bulk semiconductor can range from a
fraction of an eV up to a few eV.
We might expect that the energy dispersion relations are still parabolic in a
quantum dot. However, since only discrete energy levels can exist in a dot, each
of the original parabolic bands of the bulk case is now fragmented into an ensemble of points. The energy levels of a quantum dot can be estimated with the
particle-in-a-box model. As described in the previous paragraph (Figure 2-4), the
lowest energy for an electron in a one-dimensional potential well is
E well; 1d ¼ 1=8 h 2 =md 2
ð10Þ
Here, d is the width of the well. In a quantum dot, the charge carriers are confined in all three dimensions, and this system can be described as an infinite 3dimensional potential well. The potential energy is zero everywhere inside the well
but is infinite on its walls. We can also call this well a ‘‘box’’. The simplest shapes
of a three-dimensional box can be, for instance, a sphere or a cube. If the shape is
cubic, the Schrödinger equation can be solved independently for each of the three
translational degrees of freedom, and the overall zero-point energy is simply the
sum of the individual zero point energies for each degree of freedom [9, 60]:
E well; 3dðcubeÞ ¼ 3 E well; 1d ¼ 3=8 h 2 =md 2
ð11Þ
21
22
2 Quantum Dots
bulk semiconductor
quantum dot semiconductor
E(k)
E(k)
V(r)
y(r)
Eg(bulk)
k
Free charge carriers in a solid have
a parabolic dispersion relation ðEðkÞ z k 2 Þ.
In a semiconductor the energy bands for free
electrons and holes are separated by an energy
gap Eg . In a bulk semiconductor, the states are
quasi-continuous, and each point in the energy
bands represents an individual state. In a
quantum dot the charges are confined to a
small volume. This situation can be described
as a charge carrier confined in an infinite
Fig. 2-8
Eg(d)
r
k
d
potential well of width d. Here, the width d of
the potential well corresponds to the diameter
of the quantum dot. The only allowed states
are those whose wavefunctions vanish at the
borders of the well [7]. This leads to discrete
energy levels [7, 10]. The energy gap between
the lowest possible energy levels for electrons
and holes Eg ðdÞ is larger than that of a bulk
material Eg (bulk).
If the box is a sphere of diameter d, the Schrödinger equation can be solved by
introducing spherical coordinates and by separating the equation into a radial part
and a part that contains the angular momentum [61, 62]. The lowest energy level
(with angular momentum ¼ 0) is then
E well; 3dðsphereÞ ¼ 1=2 h 2 =md 2
ð12Þ
The effect of quantum confinement is again remarkable. More confined charge
carriers lead to a larger separation between the individual energy levels, as well as
to a greater zero-point energy. If carriers are confined to a sphere of diameter d, the
zero-point energy is higher than that for charges that are confined to a cube whose
edge length is equal to dðE well; 3dðsphereÞ > E well; 3dðcubeÞ Þ. This is because such a
sphere simply has a smaller volume ðp=6 d 3 Þ than the cube ðd 3 Þ.
An electron-hole pair can be generated in the quantum dot, for instance, by a
photoinduced process, or by charge injection. The minimum energy Eg required
for creating an electron-hole pair in a quantum dot is made up of several contributions. One contribution is the bulk band gap energy, Eg (bulk). Another
2.5 Energy Levels of a (Semiconductor) Quantum Dot
important contribution is the confinement energy for the carriers, which we call
E well ¼ E well ðE Þ þ E well ðhþ Þ. For large particles (bulk: d ! y) E well tends to zero.
We can estimate the overall confinement energy for an electron-hole pair in a
spherical quantum dot. It is the zero point energy of the potential well, or in other
words, the energy of the state of a potential box with the lowest energy. This can be
written as
E well ¼ h 2 =2m d 2
ð13Þ
where m is the reduced mass of the exciton and is given by [63]
1=m ¼ 1=m e þ 1=mh :
ð14Þ
Here m e and mh are the effective masses for electrons and holes, respectively. In
order to calculate the energy required to create an electron-hole pair, another term
(ECoul ) has to be considered. The Coulomb interaction ECoul takes into account the
mutual attraction between the electron and the hole, multiplied by a coefficient
that describes the screening of the carriers by the crystal. In contrast to E well , the
physical meaning of this term can be understood within the framework of classical
electrodynamics. However, an estimate of such a term is only possible if the wavefunctions for the electron and the hole are known. The strength of the screening
coefficient depends on the dielectric constant e of the semiconductor. An estimate
of the coulomb term yields
ECoul ¼ 1:8 e 2 =2pee 0 d
ð15Þ
This term can be quite significant, because the average distance between an electron and a hole in a quantum dot can be small [50, 64–67]. We can now estimate
the size-dependent energy gap of a spherical semiconductor quantum dot, which is
[50, 63–67]:
Eg ðdotÞ ¼ Eg ðbulkÞ þ E well þ ECoul
ð16Þ
Then, by inserting Eqs. (13) and (15) into (16) we get:
Eg ðdÞ ¼ Eg ðbulkÞ þ h 2 =2m d 2 1:8 e 2 =2pee 0 d
ð17Þ
Here we have emphasized the size dependence in each term. Equation (17) is only
a first approximation. Many effects, such as the crystal anisotropy and the spinorbit coupling, have to be considered in a more sophisticated calculation. The basic
approximation for the band gap of a quantum dot comprises two size-dependent
terms: the confinement energy, which varies as 1=d 2 , and the Coulomb attraction,
which varies as 1=d. The confinement energy is always a positive term, and thus
the energy of the lowest possible state is always raised with respect to the bulk
situation. On the other hand, the Coulomb interaction is always attractive for
23
2 Quantum Dots
4
Eg (d) [eV]
24
experiment
theory
bulk
3
2
1
4
6
8
10
d [nm]
Size dependence of the energy
gap Eg ðdÞ for colloidal CdSe quantum dots
with diameter d. The bulk value for the
energy gap is Eg ðbulkÞ ¼ 1:74 eV [63]. The
theoretical curve was obtained using Eq.
(17) with the following parameters: effective
mass of electrons/holes me ¼ 0:13m0 ,
mh ¼ 0:4m0 , m0 ¼ mass of free electrons
(m ¼ 9:1095 1031 kg) ) m ¼ 0:098 m [63];
Fig. 2-9
dielectric constant eCdSe ¼ 5:8 [176], permittivity constant e 0 ¼ 8:854 1012 C 2 N1 m2 ,
Planck’s constant h ¼ 6:63 1034 Js, 1 eV ¼
1:602 1019 J. The experimental values were
obtained in the Alivisatos group by recording
the absorption spectra of CdSe quantum dots
of different sizes and determining the size of
the quantum dots by transmission electron
microscopy (TEM).
an electron-hole pair system and therefore lowers the energy of its formation.
Because of the 1=d 2 dependence, the quantum confinement effect becomes the
predominant term for very small quantum dot sizes.
The size-dependent energy gap can be a useful tool for designing materials with
well-controlled optical properties. A much more detailed analysis on this topic
can be found, for example, in an article by Efros and Rosen [56]. Before we start
describing the physical consequences of the size-dependent band gap on the optical and electronic properties, we will first give a short overview of how quantum
dots can be fabricated in practice.
2.6
Varieties of Quantum Dots
The ultimate technique for the fabrication of quantum dots should be able to produce significant amounts of sample, with such a high control of quantum dot size,
shape and monodispersity that single-particle properties are not averaged by sample inhomogeneity. So far, ensembles of quantum dots produced by the best available techniques still show behavior deriving from a distribution of sizes, but this
field is evolving very rapidly. In this Section we give a short survey of the most
popular fabrication approaches. Different techniques lead to different typologies of
quantum dots. The confinement can be obtained in several different ways, and in
addition the quantum dot itself can have a peculiar arrangement with respect to its
2.6 Varieties of Quantum Dots
surrounding: it can be embedded into a matrix or grown onto a substrate, or it can
be a ‘‘free’’ nanoparticle. Each of these cases is strictly related to the preparative
approach chosen.
2.6.1
Lithographically Defined Quantum Dots
Lithographically defined quantum dots are formed by isolating a small region of
a two-dimensional electron system (see Section 2.4b) via tunneling barriers from
its environment. Such two-dimensional electron systems (2DES) or 2D electron
gases (2DEGs) can be found in metal-oxide-semiconductor field effect transistors
(MOSFETs) or in the so-called semiconductor heterostructures [17, 18]. Heterostructures are composed of several thin layers of different semiconductor materials
grown on top of each other, using a technique called molecular beam epitaxy
(MBE). The layer sequence can be chosen in such a way that all free charge carriers
are confined to a thin slice of the crystal, forming essentially a two-dimensional
electron system. A superstructure derived from the periodic repetition of this sequence of layers is also called a ‘‘multiple quantum well’’. One of the most widely
investigated systems is, for instance, the aluminum gallium arsenide/gallium
arsenide (AlGaAs/GaAs) quantum well. AlGaAs has the same lattice constant as
GaAs but a wider band gap, whose exact value depends on the aluminum content
of the layer. Therefore electrons in the GaAs layer are confined to this layer and
form a two-dimensional electron gas.
Quantum dot systems can be generated in a lateral or a vertical arrangement, as
shown in Figure 2-10a. In the lateral geometry, the 2DEG is locally electrostatically
depleted by applying negative voltages to electrodes deposited on top of the crystal.
We can understand this effect from the following argument. Let us assume that we
apply a negative voltage to the metal electrodes above the two-dimensional electron
gas. Because of electrostatic repulsions, electrons will be repelled by the electric
field of the electrodes, and the region of the 2DEG below the electrodes will be
depleted of electrons. A charge-depleted region behaves like an insulator. Therefore, by applying an electric field with metal electrodes of an appropriate shape it is
possible to create an island of charges insulated from the rest of the 2DEG. If the
island within the 2DEG is small enough, it behaves as a quantum dot. In the vertical geometry, a small pillar of the 2DEG is isolated by etching away the heterostructure around it. In such arrangement charge carriers are again confined in all
three dimensions.
Most of the electron transport measurements on quantum dots that have been
performed to date have used the two types of quantum dots as samples that we
have just described. The lateral arrangement offers a relatively high degree of
freedom for the design of the structure, as this is determined by the choice of the
electrode geometry. In addition, it is possible to fabricate and study ‘‘artificial molecules’’ [68–70, 71, 72–74] composed of several quantum dots linked together. With
the vertical arrangement, structures with very few electrons can be fabricated [75].
At the present time, several research efforts are focused on the investigation of
25
2 Quantum Dots
a1)
a2)
gates
metal
GaAs
AlGaAs
2DEG
one individual
quantum dot
P=O
P=O
P=
O
O=P
O=P
P=O
=O
P
P=O
P=
O
O=P
O =P
O=P
P=O
P= O
P=O
O=P
P =O
P =O
P=
O
P =O
P=
O
P =O
P=O
O= P
O=P
O= P
O=P
O
P=
P= O
O =P
P=O
P=O
P=O
=O
P=O
O= P
O=P
O
P=
O= P
O =P
O =P
P
=O
P= O
InAs
P =O
P=O
P
P=O
O=P
P =O
GaAs
O
c)
P=
b)
one individual
quantum dot
P= O
P=O
=O
O=P
O=P
O=P
O
P=
P
26
GaAs
Three different types of quantum
dots. (a1) A lithographically defined quantum
dot in lateral arrangement can be formed by
electrostatic depletion of a two-dimensional
electron gas (2DEG, shown in dark gray) via
gate electrodes. The 2DEG is formed typically
20–100 nm below the surface of a semiconductor heterostructure (usually GaAs/AlGaAs).
Application of negative voltages to metal gates
on top of the heterostructure depletes the
2DEG below the gates (shown in light gray)
and cuts out a small electron island from the
2DEG. Electrons can still tunnel onto and from
the island. Electrical contact to the 2DEG is
realized through ohmic contacts (not shown in
this picture). (a2) A vertical quantum dot can
be formed in a double barrier heterostructure.
A narrow pillar is etched out of a GaAs/
AlGaAs/GaAs/AlGaAs/GaAs heterostructure.
The AlGaAs layers (light gray) form tunnel
barriers that isolate the central GaAs region
Fig. 2-10
from the contact region. This central GaAs
region behaves now as a quantum dot (shown
in dark gray). Electrical contact is made via
metal contacts (depicted in black) on top of
the pillar and below the heterostructure.
(b) Self-assembled quantum dots: molecular
beam epitaxy (MBE) growth of InAs (dark gray)
on GaAs (light gray) first leads to the formation of an extended layer of InAs (the wetting
layer) and then to the formation of small InAs
islands. Single electrons or electron-hole pairs
(excitons) can be confined into these InAs
quantum dots, either electrically or optically.
(c) Colloidal quantum dots: these colloidal
particles, having a diameter of only a few
nanometers, are formed using wet chemistry
and can be produced for most of the II–VI,
III–V, IV–VI and some type IV semiconductors.
The surface of colloidal quantum dots is
coated with a layer of surfactant molecules that
prevents aggregation of the particles.
many-body phenomena in these quantum dot systems. Relevant examples are, for
instance, the study of the Kondo effect [76–79] and the design and the control of
coherent quantum states with the ultimate goal of quantum information processing (for references see Section 2.8).
A remarkable advantage of lithographically defined quantum dots is that their
electrical connection to the ‘‘macro-world’’ is straightforward. The manufacturing
processes are similar to those used in chip fabrication, and in principle such
structures could be embedded within conventional electronic circuits. However, as
2.6 Varieties of Quantum Dots
the geometry of these quantum dots is determined lithographically, it is limited to
the usual size and resolution limits of lithographic techniques. Even by using
electron beam lithography for the fabrication of the quantum dots, it is not possible to tailor their size with nanometer precision. Lithographically fabricated quantum dots are typically larger than 10 nm, and therefore only relatively low lateral
confining energies can be achieved.
2.6.2
Epitaxially Self-Assembled Quantum Dots
A breakthrough in the field of epitaxially grown nanostructures was the discovery
of epitaxial growth regimes that favored the formation of nanometer size islands of
semiconductor materials on suitable substrates. These islands, exhibiting quantum
dot behavior, are naturally obtained by epitaxially growing a thin layer of a low
band gap material over a higher band gap material, using MBE or MOCVD techniques [59, 80–82]. The respective crystal faces in contact must have a significant
lattice mismatch (1–8%), as in the case of InAs on GaAs [83, 84] and Ge on Si [85,
86]. During the growth, a strained film, called the ‘‘wetting layer’’, initially forms.
The maximum thickness of this layer is related to the difference between the lattice
constants of the two materials. Past this critical thickness, a 2D ! 3D transition in
the growth regime is observed, with the spontaneous formation of an array of
nanometer-sized islands (Stranski-Krastanov regime), leading to a partial release of
the strain. If the growth is not interrupted at this step, misfit dislocations form
because the energy of formation of these defects becomes smaller than the elastic
energy accumulated in the strained film. The formation of dislocations in highly
strained epilayers (when the lattice mismatch is of the order of 10% or more) before the formation of islands limits the range of possible island-substrate materials.
The shape of the islands can be controlled by the growth conditions. Usually, the
islands have a truncated pyramidal shape, but it is also possible to form, e.g., ringshaped quantum dots [87]. The final step consists in the growth, on the top of the
islands, of several layers of the substrate material, so that the dots are completely
buried and interfaces are passivated. The relative alignment of the band gaps creates a confining potential for charge carriers that accumulate inside the quantum
dots. In addition, strain fields in the proximity of the island-substrate interface,
because of the lattice mismatch between the two materials, create potentials that
modify the band gap of the quantum dots at the bottom of the island. Holes are
more likely to be localized in this region, as they are heavier than electrons.
Self-assembled quantum dots can have a diameter as small as a few nanometers,
and so very pronounced quantum size effects can be observed in these systems.
Self-assembled quantum dots have predominantly been characterized using optical
or capacitance spectroscopy in a regime where they contain only a small number
of charge carriers. Measurements on ensembles still suffer from inhomogeneous
broadening of the spectroscopic features. However, in recent years it has been
possible to look only at a few self-assembled quantum dots at a time, or even at
single dots, by reducing the number of quantum dots by mesa-etching [88] or by
27
28
2 Quantum Dots
using confocal microscopy techniques [89]. Photoluminescence from single selfassembled quantum dots is a highly efficient process, characterized by several,
narrow emission lines related to different exciton states in the dots, and is reminiscent of the emission from atoms. As mentioned already for the case of lithographically defined quantum dots, many parallels can be drawn between atoms
and quantum dots [88, 90–92]. For these reasons, quantum dots have gained the
nickname of artificial atoms. Current research efforts are devoted to the ordering
and positioning of quantum dots and also to the reduction of their size distribution. In contrast to the case of lithographically defined quantum dots, it is a challenge to make electrical contact to self-assembled quantum dots, and therefore
most of the possible applications can be found in optics. One of the major goals of
research on self-assembled quantum dots is the fabrication of non-classical light
sources from single dots. Another possible application is to use them as lightaddressable storage devices (see Section 2.7).
2.6.3
Colloidal Quantum Dots
Colloidal quantum dots are remarkably different from quantum dots formed by the
systems mentioned above, as they are chemically synthesized using wet chemistry
and are free-standing nanoparticles or nanocrystals grown in solution [51]. Colloidal quantum dots are just a subgroup of a broader class of materials that can be
synthesized at the nanoscale level using wet chemical methods. In the fabrication
of colloidal nanocrystals, the reaction chamber is a reactor containing a liquid
mixture of compounds that control the nucleation and the growth. In a general
synthesis of quantum dots in solution, each of the atomic species that will be part
of the nanocrystals is introduced into the reactor in the form of a precursor. A
precursor is a molecule or a complex containing one or more atomic species
required for growing the nanocrystals. Once the precursors are introduced into the
reaction flask they decompose, forming new reactive species (the monomers) that
will cause the nucleation and growth of the nanocrystals. The energy required to
decompose the precursors is provided by the liquid in the reactor, either by thermal
collisions or by a chemical reaction between the liquid medium and the precursors, or by a combination of these two mechanisms [93].
The key parameter in the controlled growth of colloidal nanocrystals is the presence of one or more molecular species in the reactor, here broadly termed ‘‘surfactants’’. A surfactant is a molecule that is dynamically adsorbed to the surface of
the growing quantum dot under the reaction conditions. It must be mobile enough
to provide access for the addition of monomer units, while stable enough to prevent the aggregation of nanocrystals. The choice of surfactants varies from case to
case: a molecule that binds too strongly to the surface of the quantum dot is not
suitable, as it would not allow the nanocrystal to grow. On the other hand, a weakly
coordinating molecule would yield large particles, or aggregates [94]. Some examples of a suitable surfactant include, for instance, alkyl thiols, phosphines, phosphine oxides, phosphates, phosphonates, amides or amines, carboxylic acids, and
2.7 Optical Properties of Quantum Dots
nitrogen-containing aromatics. If the growth of nanocrystals is carried out at high
temperatures (for instance at 200–400 C), then the surfactant molecules must be
stable under such conditions in order to be a suitable candidate for controlling the
growth.
At low temperatures or, more generally, when the growth is stopped, the surfactants are more strongly bound to the surface of the nanocrystals and provide their
solubility in a wide range of solvents. This coating allows for great synthetic flexibility in that it can be exchanged with another coating of organic molecules having
different functional groups or polarity. In addition, the surfactants can be temporarily removed and an epitaxial layer of another material with different electronic,
optical, or magnetic properties can be grown on the initial nanocrystal [95, 96].
By controlling the mixture of surfactant molecules that are present during the
generation and the time of growth of the quantum dots, excellent control of their
size and shape is possible [93, 97, 98].
Since colloidal nanocrystals are dispersed in solution, they are not bound to any
solid support as is the case for the other two quantum dots systems described
above. Therefore, they can be produced in large quantities in a reaction flask and
can later be transferred to any desired substrate or object. It is, for example, possible to coat their surface with biological molecules such as proteins or oligonucleotides. Many biological molecules perform tasks of molecular recognition with an
extremely high efficiency. This means that ligand molecules bind with very high
specificity to certain receptor molecules, similarly to a key-lock system. If a colloidal quantum dot is tagged with ligand molecules, it specifically binds to all the
positions where a receptor molecule is present. In this way it has, for example,
been possible to make small groupings of colloidal quantum dots mediated by
molecular recognition [99–101] and to label specific compartments of a cell with
different types of quantum dots [102–104].
Although colloidal quantum dots are rather hard to connect electrically, a few
electron transport experiments have been reported. In these experiments, nanocrystals were used as the active material in devices that behave as single-electron
transistors [105, 106] (see also Section 2.8).
2.7
Optical Properties of Quantum Dots
2.7.1
Absorption and Emission Spectra
In chemistry, size-dependent optical properties of colloidal semiconductor particles have been observed since the beginning of the last century, for instance in
CdS colloids [107], although it was only toward its end [65] that this fact was
rationalized in terms of ‘‘size quantization’’. As we have discussed in Section 2.5,
the most striking effect in semiconductor nanoparticles is the widening of the gap
Eg between the highest occupied electronic states (the top of the original valence
29
30
2 Quantum Dots
absorption /
emission
1.4
1.2
1.0
0.8
0.6
0.4
0.2
0.0
450
500
550
600
650
700
750
wavelength [nm]
Absorption (plain lines) and
emission spectra (dotted lines) of colloidal
CdSe quantum dots of different sizes. The
absorption peaks of green/yellow/orange/red
Fig. 2-11
fluorescent nanocrystals of 2.3/4.0/3.8/4.6 nm
diameter are at 507/547/580/605 nm. The
fluorescence peaks are at 528/57/592/637 nm.
band) and the lowest unoccupied states (the bottom of the original conduction
band) [108]. This directly affects the optical properties of quantum dots as compared to the corresponding bulk material. The minimum energy needed to create
an electron-hole pair in a quantum dot (an ‘‘exciton’’) is defined by its band gap
(Eg ). Light with energy lower than Eg cannot be absorbed by the quantum dot.
Since the band gap depends on the size of a quantum dot, the onset of absorption
is also size dependent [50, 93]. In Figure 2-11 we show that smaller quantum dots
have an absorption spectrum that is shifted to shorter wavelengths with respect to
larger quantum dots and to the bulk material.
Excitons in semiconductors have a finite lifetime because of the recombination
of the photo-excited electron-hole pair. In quantum dots, the energy released upon
exciton annihilation is too large to be dissipated by vibrational modes. Instead it
is released in the form of emitted photons. Radiative decay through emission of
photons, in other words fluorescence, is a highly probable decay channel in quantum dots [109].
As is the case of organic fluorophores, the range of energies emitted from a colloidal dot sample after excitation is centered at a value that is smaller than the one
required to excite the sample (and which must be at least as large as its band gap).
In other words, the wavelength of the fluorescence is longer than that of the
absorbed light. The shift between the lowest energy peak in the absorption spectrum of a quantum dots and the corresponding emission peak is called the ‘‘Stokes
shift’’ (Figure 2-11) [9]. In order to explain the Stokes shift, a more advanced
model than that presented in Figure 2-8 is necessary.
In quantum dots, the Stokes shift is explained by examining their complex exciton structure [110–113]. More complex theoretical models and calculations show
that the ground state of an exciton in a dot has a total angular momentum equal to
2.7 Optical Properties of Quantum Dots
Fig. 2-12 Colloidal CdSe quantum dots of
different size dissolved in chloroform. The size
of the quantum dots increases from the left to
the right vial. (a) Photograph of the solutions.
(b) Photograph of the solutions upon UVillumination from below. The different colors of
fluorescence can be seen.
zero. In the dipole approximation, the creation of an exciton through absorption of
a photon leads to an exciton state having angular momentum equal to G1. The
energy needed for this excitation is the absorption energy. In quantum dots, this
excited state now relaxes very fast to a state with an angular momentum of 2. This
relaxation is non-radiative. In first order, this state cannot relax to the ground state
with angular momentum 0 by emitting a photon, because only transitions that
change the angular momentum by G1 are allowed. Since no photon can be
emitted in first order, this state is called ‘‘dark exciton’’. However, slight perturbations of the crystal lattice and even a weak coupling with phonons allow this state
to relax under emission of a photon. As a result, the decay time of the fluorescence
is long, and the fluorescence energy is red-shifted with respect to the absorption
band edge energy [114]. The dark exciton model is supported by many experimental data, in particular by magnetic field-dependent lifetime measurements.
The position of the luminescence peak is also dependent on the average quantum
dot size, and its width is correlated to the nanocrystals size distribution (Figures
2-11 and 2-12). Consequently, the maximum of the emission spectrum and its
width can be used to estimate the mean size and the size distribution during
nanocrystal growth.
2.7.2
Spectral Diffusion and Blinking
Several crucial differences appear between self-assembled and colloidal quantum
dots. For instance, stable, ultra-narrow fluorescence peaks have been observed in
31
32
2 Quantum Dots
the emission spectra of single, self-assembled quantum dots. Colloidal semiconductor quantum dots can have a very narrow size distribution, as observed under a
transmission electron microscope. Nevertheless, in ensemble measurements their
emission spectra still have a full width at half maximum of several meV [115, 116].
Although this broad range of emission energies was initially ascribed to the residual size distribution, it is now ascribed to an intrinsic property of colloidal quantum dots. Although the emission peak from a single colloidal quantum dot can
be less than 0.1 meV wide, but its emission energy shifts randomly over time. This
behavior is referred to as ‘‘spectral jumps’’ or ‘‘spectral diffusion’’ [116]. First observed by Empedocles et al. at cryogenic temperatures, spectral diffusion now can
also be observed at room temperature [117, 118]. Spectral diffusion is likely related
to the local environment of the quantum dots, which creates rapidly fluctuating
electric fields that may perturb the energy levels of the system. Similarly, spectral
diffusion can also be observed in single organic fluorophores [119]. Conversely,
self-assembled quantum dots embedded in a matrix do not exhibit spectral jumps,
because their local environment does not change with time.
In self-assembled quantum dots, multiexciton states can be observed and studied
at high pumping power, but they have never been observed in colloidal quantum
dots. The absence of multiexcitons in single colloidal quantum dots is believed
to be correlated with the fluorescence intermittence observed in these systems
[120]. The fluorescence emission from a single quantum dot exhibits a dramatic
on/off behavior, referred to as ‘‘blinking’’ and is another spectroscopic feature that
distinguishes colloidal from self-assembled quantum dots [121, 122]. This behavior
is similar to that of single organic fluorophores [123]. Blinking in CdSe quantum
dots can readily be observed in a standard epifluorescence microscope. In a nanocrystal, off time periods may vary from milliseconds to several minutes. The
mechanism responsible for the blinking is believed to involve an Auger photoionization of the colloidal quantum dots. In this model, if two electron-hole pairs
are present simultaneously inside the quantum dots, the energy released by the
annihilation of one electron-hole pair is transferred to the other pair. This excess of
energy may eject one of the carriers into the environment of the quantum dot,
leaving the quantum dot charged. If an electron-hole pair is created during this
time, the energy released by its recombination is transferred to the third remaining
carrier via a non-radiative process. This is why ionized quantum dots do not emit.
If the ejected carrier returns inside the quantum dot or if the quantum dot is neutralized, radiative emission is restored. The probability of occurrence of Auger
processes in nanocrystals is higher than in the bulk because of the breakdown
of translational symmetry. The probability is also related to the spatial overlap of
the carrier wavefunctions, and for this reason is higher in colloidal than in selfassembled quantum dots, the former being usually much smaller than the latter.
The probability of Auger processes increases further in the case of finite and
defective barriers, which offer a wide range of electronic states where the excited
carriers can be localized. Again, this is the case with colloidal quantum dots, but
not with self-assembled quantum dots, which are buried in an inorganic, defectfree, thick matrix. A more detailed description of the Auger effect in colloidal
quantum dots can be found in the literature [116, 124, 125].
2.7 Optical Properties of Quantum Dots
absorption
1.4
1.2
1.0
0.8
0.6
0.4
0.2
0.0
200
400
600
800
1000
wavelength [nm]
Fig. 2-13
Absorption spectrum of colloidal gold nanocrystals of 10 nm diameter.
2.7.3
Metal Nanoparticles
Colloidal metal nanoparticles can have optical absorption spectra with an absorption peak that looks similar to the absorption peak of colloidal semiconductor particles (compare, for instance, Figures 2-11 and 2-13). However, this absorption does
not derive from transitions between quantized energy states. Instead, in metal
particles, collective modes of motion of the electron gas can be excited. They are
referred to as surface plasmons [126, 127]. The peak in the absorption spectrum is
the resonance frequency for the generation of surface plasmons. The size dependence of the plasmon frequency is negligible. For instance, the absorption maximum for colloidal gold nanocrystals hardly shifts for nanoparticles in the size
range between 5 and 30 nm. This contrasts with the dramatic size effect in the
absorption of colloidal semiconductor nanocrystals when the diameter is changed
by only fractions of a nanometer (see Figure 2-11).
2.7.4
Overview of Some Selected Applications
Many applications of the quantum-mechanical aspects of quantum dots can be
found in optics. As in the more general case of atoms or molecules, quantum dots
can be excited either optically or electrically. Regardless of the nature of the excitation, quantum dots may emit photons when they relax from an excited state to
the ground state. Based on these properties, quantum dots may be used as lasing
media, as single photon sources, as optically addressable charge storage devices, or
as fluorescent labels.
Self-assembled quantum dots incorporated into the active layer of a quantum
well laser significantly improve the operation characteristics of the laser. This is
because of the zero-dimensional density of states of quantum dots. In quantum
dot lasers, the threshold current density is reduced, its temperature stability is
improved, and the differential gain is increased. A first demonstration of the lasing
33
34
2 Quantum Dots
operation of a quantum dot laser structure was given as early as 1994 [128, 129].
Since then, the lasing characteristics have been improved by better control of the
growth of the self-organized quantum dot layers. Today, quantum dot lasers are on
the road toward commercialization [130]. Optical gain and stimulated emission
have been observed also from CdSe and CdS colloidal nanocrystal quantum dots
[131]. On the basis of this result, it is conceivable that optical devices could also be
built using self-assembly of colloidal quantum dots.
Quantum dots have been used not only as conventional laser sources but also
as ‘‘non-classical’’ light sources. Photons emitted from thermal light sources have
characteristic statistical correlations: in arrival time measurements it is found that
they tend to ‘‘bunch’’ together (super-poissonian counting statistics). For applications in quantum information processing, however, it would be desirable to emit
single photons one at a time (sub-poissonian counting statistics, or anti-bunching).
In the last few years it has been possible to demonstrate first prototypes of such
single-photon sources based on single quantum dots [132–136].
Self-assembled quantum dots have also been discussed as the basis of an alloptical storage device. In such a device, excitons are optically generated and electrons and holes are stored separately in coupled quantum dot pairs [137]. Applying
an electric field, electron and hole can be forced to recombine and generate a photon which provides an optical read-out.
Colloidal quantum dots have also been used for the development of lightemitting diodes [138, 139]. In these devices, colloidal quantum dots are incorporated into a thin film of conducting polymer. Furthermore, colloidal quantum dots
have been used for the fabrication of photovoltaic devices [140, 141].
Chemically synthesized quantum dots fluoresce in the visible range with a
wavelength tunable by the size of the colloids. The possibility of controlling the
onset of absorption and the color of fluorescence by tailoring the size of colloidal
quantum dots makes them interesting objects for the labeling of biological structures [102, 103] as a new class of fluorescent markers. The tunability combined
with extremely reduced photobleaching makes colloidal quantum dots an interesting alternative to conventional fluorescent molecules [51]. Possible biological applications of fluorescent colloidal quantum dots will be discussed in detail in the
other chapters of this book.
2.8
Some (Electrical) Transport Properties of Quantum Dots
Electron transport through ultrasmall structures such as quantum dots is governed
by charge and energy quantization effects. Charge quantization comes into play for
structures with an extremely small capacitance. The capacitance of a nanostructure
– which, roughly speaking, is proportional to its typical linear dimension – may
become so small that the energy required to charge the structure with one additional charge carrier (electron, hole, cooper pair) exceeds the thermal energy available. In this case, charge transport through the structure is blocked – an effect
2.8 Some (Electrical) Transport Properties of Quantum Dots
which has appropriately been termed the ‘‘Coulomb blockade’’ (CB) effect. As will
be shown below, this effect can be exploited to manipulate single electrons within
nanostructures. Also because of the smallness of the structures, energy quantization may become considerable. In contrast to bulk structures, charge carriers
within a quantum dot are only allowed to occupy discrete energy levels, as in the
case of electrons within an atom. In a manner resembling scattering effects known
from atomic and nuclear physics, the occurrence of these discrete energy levels can
modify the charge transport characteristics of ultrasmall devices. In metallic nanostructures, Coulomb blockade can also occur without ‘‘quantum aspects’’, as the
energy level spacing in these structures is usually too small to be observable. In
fact, energy quantization is the reason for the name ‘‘quantum dot’’ and differentiates this from most of the metallic nanostructures. However, in extremely
small metallic nanoparticles and clusters, quantum effects can also be considerable. In this section we give a short introduction of the history and fundamentals
of the Coulomb blockade effect. We will discuss single-electron tunneling transport
through quantum dots and we will also give an overview of the possible applications of these effects.
2.8.1
Coulomb Blockade: Basic Theory and Historical Sketch
The Coulomb blockade effect was originally observed in experiments on small
metallic or superconducting particles [142–145]. In these experiments, nanometersized metallic grains were embedded within a metal-oxide-metal tunnel junction.
Electronic measurements performed on these systems at low temperatures (T A
1 K) revealed anomalous behavior of the resistance (or differential capacitance) at
zero bias. It was realized that this behavior was caused by the extremely small
capacitance of the metallic particles. In a simple model, a spherical particle of
diameter d embedded in a dielectric with permittivity e has a capacitance
ð18Þ
C ¼ 2pee 0 d
and therefore at low enough temperatures the charging energy3)
EC ¼
e2
e2
¼
2C 4pee 0 d
ð19Þ
required for the addition of a single electron to the particle may well exceed the
thermal energy kB T. As an example, a metallic particle of radius 10 nm embedded
3) Although the charging energy EC ¼ e 2 =4pee 0 d
(Eq. (19)) looks very similar to the Coulomb
energy ECoul ¼ 1:8 e 2 =2pee 0 d (Eq. (1)), the
two terms describe different situations. The
charging energy describes the energy that is
needed to add an additional electron to an
already negatively charged system (a sphere
with capacity C), and therefore it is repulsive
ðE > 0Þ. The Coulomb energy describes the
attraction ðE < 0Þ between an electron and a
hole within a sphere.
35
36
2 Quantum Dots
in a dielectric with e ¼ 4 has a charging energy of EC ¼ 18 meV, which corresponds to a temperature of about 200 K. Below this temperature (as a rule of
thumb, EC > 10kB T should be satisfied), it is impossible to add or remove a single
charge to the particle at zero bias, and therefore electronic transport is blocked.
This explains the highly increased resistance of the tunnel junctions at low temperatures observed in the early experiments. However, the Coulomb blockade can
be lifted if enough energy is supplied by applying a bias over the structure. It is
found that for jVbias j > e=2C the conductance starts to rise from its suppressed
value. The voltage interval ½e=2C; þe=2C over which conduction is suppressed is
often referred to as the ‘‘Coulomb gap’’. Apart from the smallness of the conductive island, the other essential requirement for Coulomb blockade to be observed is
that the metallic island or particle be isolated from the contacts via tunnel barriers.
This ensures that the charge on the island is ‘‘sufficiently well quantized’’. If the
island were coupled more strongly to the contacts, charge could ‘‘leak out’’ and
thus destroy the effect. To be more precise, the energy uncertainty of the electrons
due to the finite lifetime t of its energy state (given roughly by the RC time of the
island, where R is resistance) should be less than the charging energy, i.e.
DE A
p
p
e2
¼
<
t RC C
ð20Þ
Solving for R yields as a condition for the resistance R > p=e 2 which just means
that R should be of the order of the resistance quantum4).
Many years after the initial experiments, the Coulomb blockade effect was
observed for the first time in a lithographically defined metallic structure [146]. In
contrast to the earlier experiments, the effect could now be observed for a single
metallic island rather than averaged over a large population of nanoparticles.
Shortly after, Coulomb blockade was also observed in a semiconductor microstructure [147]. The use of lithographic techniques facilitates the fabrication of
well-defined three-terminal geometries. In such a geometry, in addition to the
source and drain contacts, a gate electrode is placed near the electron island. With
this gate the electrostatic energy of the island can be tuned, and therefore the
number of electrons on the island can be changed one by one. Variation of the gate
voltage therefore leads to almost periodic conductance oscillations, which indicate
the charging of the island with single electrons (see discussion below).
In most cases, for a metallic nanostructure, Coulomb blockade is a purely classical (electrostatic) phenomenon in the sense that the energy level spacing DE
in the structure is much smaller than the thermal energy5) kB T which is due to
the metal’s high density of states at the Fermi level. This is strikingly different
from the case of semiconductor nanostructures, for which the level spacing can
become comparable to or even larger than the charging energy. Hence, in semiconductor nanostructures the energy required to add or remove single electrons is
strongly modified by the quantum-mechanical energy levels. For this reason, the
4) h=e 2 is the ‘‘resistance quantum’’. It is equal
to the resistance of a single non-degenerate
conduction mode of a one-dimensional
conducting channel connecting two large
reservoirs.
5) kB is the Boltzmann constant
2.8 Some (Electrical) Transport Properties of Quantum Dots
term ‘‘quantum dots’’ is usually used for such ultrasmall semiconductor structures. Following the initial experiments mentioned above, a huge body of work has
been published on transport properties of metallic islands and semiconductor
quantum dots in the Coulomb blockade regime.
2.8.2
Single-Electron Tunneling
A number of excellent reviews of the theoretical and experimental aspects of Coulomb blockade and single-electron tunneling have been given in the last decade.
An early review on single-electron tunneling in metallic islands has been given by
Averin and Likharev [148]. An updated treatment by G. Schön emphasizing other
aspects can be found in [149]. Theoretical aspects of Coulomb blockade in semiconductor nanostructures are surveyed in [150]. A good review of transport experiments on semiconductor quantum dots is in [53]. We also recommend the shorter
reviews in [151] and [152]. In this Section we give a brief introduction to a simple
model of a single-electron transistor which already captures many of the relevant
aspects. For detailed discussions we refer to the more specialized reviews and references therein.
A schematic representation of a single-electron transistor is given in Figure 2-14.
It consists of a small conducting island connected to its surroundings via two tunneling barriers and a gate electrode. The capacitances of the barriers and the gate
are denoted CS , CD and CG , and the drain-source and gate voltages V ¼ VD VS
and VG . The total capacitance of the island is given by C ¼ CS þ CD þ CG. The
electrostatic energy of the island in this model is given by
EðN; Q G Þ ¼
ðNe Q G Þ 2
2C
Fig. 2-14 Schematic circuit diagram for a
single-electron transistor. The electron island is
indicated by the black dot. It is connected to
source and drain contacts via tunneling
barriers having capacitances CS and CD .
Additionally, the electrostatic energy of the
ð21Þ
island can be tuned with a capacitively coupled
gate (capacitance CG ). In this circuit, the
source contact has been set to ground and the
gate voltage is applied with respect to the
ground potential.
37
38
2 Quantum Dots
where the integer N is the number of electrons on the island, e is the electronic
charge and the gate charge Q G ¼ CD VD þ CG VG þ CS VS (in our circuit diagram
we set VS ¼ 0 as the reference potential). Terms independent of N have been
omitted in this expression. EðN; Q G Þ defines a set of parabolas with minima at
Ne ¼ Q G . Whereas the gate charge Q G can be varied continuously by the external
voltage sources, in the Coulomb blockade regime the charge on the island can only
vary in integer numbers. For different gate charges (gate voltages) the island may
therefore be occupied by a different number of electrons. In other words, the gate
voltage can be used to tune the number of electrons on the island. The charge can
fluctuate, however, if the energies for two successive occupation numbers are
degenerate, i.e., if EðN þ 1; Q G Þ ¼ EðN; Q G Þ. At these points the Coulomb blockade is lifted and charges can be added to or removed from the dot, i.e., the conductance of the dot becomes finite. Solving
for the gate charge leads to the condition for charge fluctuation6) Q G ðNÞ ¼ N þ 12 e. The distance between two adjacent
degeneracy points is therefore Q G ðNþ1Þ Q G ðNÞ ¼ e. Let us consider the special
case VDS ¼ 0 (zero bias). Then, Q G ¼ CG VG and DVG ¼ V G ðNþ1Þ V G ðNÞ ¼ e=CG .
This shows that the gate voltages at which a finite conductance can be measured
are equally spaced – leading to the periodic conductance oscillations of a singleelectron transistor in the Coulomb blockade regime.
In a thermodynamic description of the system, the charge on the island can
fluctuate if the probability to find N electrons on the island equals that to find
N þ 1 electrons, or PðNÞ ¼ PðN þ 1Þ, where the equilibrium probability P is given
by
PðNÞ ¼
1
expððEðNÞ mNÞ=kB TÞ
Z
ð22Þ
Z is the partition function (Zustandssumme) [153] and m is the chemical potential
of the contacts. The condition for fluctuation thus becomes EðN þ 1Þ EðNÞ ¼ m.
The left side of this equation is called the addition energy for the ðN þ 1Þ-th electron. If a finite bias V is applied over the system, the electrochemical potentials of
the two reservoirs differ by eV. In this case a ‘‘conductance window’’ opens, and
Coulomb blockade is lifted for
m ewV < EðN þ 1Þ EðNÞ < m þ eð1 wÞV
ð23Þ
where w ð0 a w a 1Þ is the fractional voltage drop over the drain barrier.
6) by using Eq. (19) the condition for charge
fluctuation EðN þ 1; Q G Þ ¼ EðN; Q G Þ
becomes
2
2EQ G þ Q G 2
2
ððN þ 1ÞE Q G Þ =2C ¼ ðNE Q G Þ =2C
2
, N 2 E 2 þ 2NE 2 þ E 2 2NEQ G
2
, ðN þ 1Þ E 2ðN þ 1ÞEQ G
þ Q G 2 ¼ N 2 E 2 2NEQ G þ Q G 2
¼ N 2 E 2 2NEQ G þ Q G 2
, 2NE 2 þ E 2 2EQ G ¼ 0
, 2EQ G ¼ ð2N þ 1ÞE 2
, Q G ¼ N þ 12 E
2.8 Some (Electrical) Transport Properties of Quantum Dots
Fig. 2-15 Schematic depiction of electrostatic
energy E/EC, occupation number N and conductance G of the single-electron transistor
(SET) shown in Figure 2-14 as a function of
the gate charge QG . The occupation of the
SET changes when the energy for N electrons
equals that of N þ 1. At the points where
the charge can fluctuate, the conductance
becomes finite and the Coulomb blockade
is lifted. In real systems, the line shape of
the conductance peaks is determined by
temperature and/or lifetime broadening.
I–V curves can be calculated from a master equation for the occupation
probability of the Coulomb blockade island. For the single-electron transistor with
two tunneling barriers and one island, the master equation is [148, 149]
d
pðN; tÞ ¼ ½GL!I ðN 1Þ þ GR!I ðN 1Þ pðN 1; tÞ
dt
þ ½GI!L ðN þ 1Þ þ GI!R ðN þ 1Þ pðN þ 1; tÞ
½GL!I ðNÞ þ GI!R ðNÞ þ GR!I ðNÞ þ GI!R ðNÞ pðN; tÞ
ð24Þ
where pðN; tÞ is the nonequilibrium probability of finding N electrons on the
island at time t. The GX!Y ðNÞ are tunneling rates for transitions from X to Y,
where X and Y can be L (left contact), R (right contact) or I (island). An example of
a typical IV curve obtained from the master equation is shown in Figure 2-16. In
the Coulomb blockade regime (when the gate charge
is QG ¼ Ne), the current in
the Coulomb gap around V ¼ 0 is zero. For Q G ¼ N þ 12 E, the Coulomb blockade is lifted and the IV curve is linear with finite slope at V ¼ 0.
For a quantum dot, one also has to account for the energy quantization of the
electrons on the island. In the simplest model, sometimes referred to as ‘‘constant
interaction (CI) model’’, one assumes that the electrons successively occupy single-
39
2 Quantum Dots
2
1
I
40
0
-1
-2
-2
0
2
V
Calculated I–V-curves for a single
electron transistor for two different values of
the gate charge. The source and drain tunnel
resistances and capacitances are assumed
equal, the charging energy is 100 kB T. For gate
charge QG ¼ 0 the SET is in the Coulomb
blockade regime. The I–V curve exhibits the
‘‘Coulomb gap’’. For a gate charge QG ¼ 0:5
Fig. 2-16
the Coulomb blockade is lifted (cf. Fig. 2-15),
and the I–V curve is linear around V ¼ 0. The
calculation was performed using the program
SETTRANS [SETTRANS by Alexander N.
Korotkov7) is available for noncommercial use
under http://hana.physics.sunysb.edu/set/
software/index.html].
particle energy levels ei (see Figure 2-48)), and the mutual electrostatic repulsion of
the electrons is accounted for by a classical capacitance. Thus the energy for N
electrons on the quantum dot is
EðNÞ A
N
ðNeÞ 2 X
þ
ei
2C
1
ð25Þ
The addition energy therefore becomes
EðN þ 1Þ EðNÞ ¼
ð2N þ 1Þe 2
þ eNþ1
2C
ð26Þ
and the distance between two adjacent conductance resonances follows as
DmN ¼ EðN þ 1Þ 2EðNÞ þ EðN 1Þ ¼
7) The program SETTRANS by Alexander N.
Korotkov is available for noncommercial use
under http://hana.physics.sunysb.edu/set/
software/index.html
8) In order to avoid confusion, we here denote
the discrete energy levels of electrons in a
e2
þ DeN
C
ð27Þ
potential well by ei. These energy levels are
identical to those shown in Figure 2-4. For a
three-dimensional potential box we obtain
ei ¼ i 2 h 2 =2md 2 , i ¼ 1; 2; 3; . . . (cf. Eq. (12)).
The lowest possible energy ði ¼ 1Þ is
h 2 =2md 2 .
2.8 Some (Electrical) Transport Properties of Quantum Dots
where DeN ¼ eNþ1 eN is the Nth level spacing. In the CI model, the distance
between adjacent conductance maxima is given by a classical charging term and
the spacing between single particle energy levels. This shows how the quantum
nature of quantum dots directly enters its electrical properties.
In systems where the CI model can be applied, the spectrum of the quantum dot
can, in principle, be deduced from transport or capacitance measurements, experimental techniques which therefore have been termed transport or capacitance
spectroscopy. The CI model is quite reasonable for systems with very few electrons,
but in many situations many-body effects lead to strong deviations from this simple behavior. The relative importance of energy quantization versus charging energy can be judged from the magnitude of the terms e 2 =C and De in the expression
above. Typically, the capacitance C varies with the diameter d of a system like
C A 2pee 0 d (cf. Eq. (18)), and thus EC z e 2 =ee 0 d (cf. Eq. (19)), whereas the quantization energies are of the order of De A h 2 =2m d 2 (cf. Eq. (13)). Most trivially, for
smaller systems quantum effects become stronger.
2.8.3
Tunneling Transport – Line Shape of Conductance Peaks
A slightly different approach to treating electron transport through quantum dots
considers conduction as a transmission problem through an electron island. In
this formulation, single-electron tunneling through quantum dots bears similarities to the well-known description of resonant tunneling transport. Resonant tunneling occurs when a system with discrete energy levels is connected to two reservoirs via tunneling barriers. In the case of quantum dots, the discreteness of the
energies naturally comes in because of either Coulomb blockade or quantum confinement, or both. In Figure 2-17, such a system is depicted schematically. The
Fig. 2-17 Electron transport through a
quantum dot viewed as double barrier
tunneling. The ‘‘resonant energies’’ of the
quantum dot are given by the addition
energies discussed in the text. In the constant
interaction model the difference between
adjacent resonances is approximated by a
charging term and the single particle energy
level spacing. The chemical potentials of the
reservoirs are denoted mL and mR , the tunneling
rates through the barriers are GL and GR .
41
42
2 Quantum Dots
tunneling rates through the left and right tunnel barriers are denoted by GL and
GR ; the quantum level in the dot has energy e1. One can show that the transmission at energy E through such a system is approximately given by
TðEÞ ¼
GL GR
GL þ GR 2
ðE e1 Þ 2 þ
2
ð28Þ
where the tunneling rates have been assumed to be energy independent. The current is then given by
I¼
ð
2e
½ f ðEÞ f ðE eVds ÞTðEÞ dE
h
ð29Þ
Here, f ðEÞ denotes the Fermi function, and the factor of two accounts for spin
degeneracy. From this expression, the conductance of the system can be determined and hence the shape of the conductance oscillations in the Coulomb blockade regime. Depending on the relative size of the coupling energies hGL ; hGR and
the thermal energy kB T, the line shape is dominated by lifetime broadening of the
energy state on the island or by temperature broadening. In the low bias regime
ðeVDS A 0Þ, the Fermi function in the above expression can be expanded, which
yields for the conductance
G¼
2e 2
h
ð
f 0 ðEÞTðEÞ dE
ð30Þ
Thus, if temperature broadening dominates, the conductance oscillations have
the shape of the derivative of the Fermi function (z cosh2 ðE=2kB T)). If lifetime
broadening dominates, the conductance peaks are of Lorentzian shape.
2.8.4
Some Applications
The peculiar properties of quantum dots described in the previous sections may be
used in a variety of device applications. The most important of these features are
the discrete energy spectrum, the smallness of the capacitance, and (which is
related) the possibility of manipulating single charges. The survey of device applications given in the following sections is by no means complete. It is intended to
demonstrate the wealth of work done in the field and the many directions pursued.
The possibility of manipulating single charges in single-electron tunneling devices led to early proposals for metrological applications of the Coulomb blockade
effect. Devices called ‘‘electron turnstiles’’ can be used to count single electrons
and could form the basis of accurate current standards based on the definition of
the second. Electron turnstiles basically consist of a number of electron islands
2.8 Some (Electrical) Transport Properties of Quantum Dots
connected together via tunneling barriers. The potentials of the islands can be
controlled individually using gates. If the gate voltages are changed in a particular
operation cycle, the potentials of the islands can be changed in such a manner that
single electrons are ‘‘pumped’’ through the system. If the operation cycle is repeated at a frequency f , the resulting current is I ¼ ef . Thus, the current is directly
linked to the definition of the second via frequency [154–156]. The same principles
can be used to construct capacitance standards [157].
Single-electron transistors were also discussed as the basis of logic circuits.
Some concepts try to implement conventional CMOS (complementary metal-oxidesemiconductor) logic using single-electron transistors as replacements for n-MOS
and p-MOS transistors (depending on whether the SET is in the CB region or
whether CB is lifted) [158, 159]. In other concepts, the logical states are essentially
represented by single electrons, e.g., in quantum dot cellular automata [160]. All of
these proposals face strong conceptual and technical challenges. So far, it does
seem unlikely that single electronics offers a competitive alternative to traditional
CMOS-based logic. Among the problems are the required smallness of the quantum dots for reliable room temperature operation, the sensitivity to background
charges, the need for strategies for reproducible placing and interconnecting of
quantum dots to build VLSI structures, and the lack of gain of single-electron
transistors [161].
Because of their low capacitance, nanometer-sized metallic islands or quantum
dots are extremely sensitive to neighboring charges. Whereas in some situations
this is unfavorable, e.g., if fluctuating background charges disturb the operation
of an SET, this effect can also be exploited for device applications. A promising
application of the charge sensitivity of quantum dots is the fabrication of singleelectron memories. In one approach, the floating gate memory technique which
is already used in CMOS memories is adapted to utilize single-electron effects. A
single-electron transistor is used as an electrometer which senses the charging or
discharging of a floating gate [162]. The charge state of the floating node can thus
be used as a memory whose state is read out by an SET.
Some of the most promising applications of single-electron devices lie in the
area of sensorics. An obvious sensor application is charge sensing and electrometry: SETs are the most sensitive charge detectors
to date with an equivalent
pffiffiffiffiffiffiffi
charge noise of the order of 103 to 105 E= Hz at 10 Hz operation
pffiffiffiffiffiffiffifrequency
(see, for example, [163, 164]). A charge noise of about 2 105 E= Hz could be
5
achieved
pffiffiffiffiffiffiffi at a higher operation frequency of 4 kHz [165] and even 1:2 10
E= Hz at 1.1 MHz (with an ‘‘RF-SET’’) [166]. SET-based electrometers were
employed to characterize other physical systems, e.g., for the characterization of
quantum Hall edge states [167] or for capacitance spectroscopy of vertical quantum
dots [168]. SET electrometers were even employed as charge sensors for scanning
probe microscopy (SPM) [169]. An interesting application not explicitly based on
charge sensitivity is the utilization of quantum dots for the detection of single
photons in the far-infrared range [170]. More generally, single-electron transistors
are expected to prove useful in many low-noise analog applications, in particular in
the amplification of quantum signals [171].
43
44
2 Quantum Dots
Most of the device proposals mentioned above do not rely on the quantum
nature of quantum dots, but simply make use of the Coulomb blockade effect
based on the ultrasmall capacitance of small electron islands. The majority of
applications of energy quantization within quantum dots are naturally found in
optics (see Section 2.7). But also in electronics, the possible utilization of quantum
dot energy levels for quantum information processing has recently attracted considerable attention. To build a quantum computer, techniques have to be developed
to produce and manipulate ‘‘qubits’’ (quantum bits) – the information unit in
quantum information processing and quantum computing. Qubits can be realized
in any physical system with two quantum states which can be identified as logical
‘‘0’’ or ‘‘1’’ (in Dirac notation j0i and j1i). The difference between classical and
quantum logical states is that a quantum system can also be in a superposition of
j0i and j1i [172]. Certain classes of computational problems – e.g. factorization of
large numbers – are expected to be solved more efficiently on quantum computers.
Possible realizations of qubits are the spin states of particles (spin up – spin down),
the polarization of photons (vertical – horizontal) or other physical systems which
can be approximately described as two-level quantum systems. Implementations of
quantum computers in quantum dots make use of the quasi-molecular states
formed in two coupled quantum dots [173]. The j0i and j1i states correspond to
an electron localized on one or other of the two dots. Superpositions can be formed
by strongly coupling the two dots together. First experimental realizations of such
artificial molecular states in quantum dots have already been reported [174, 175].
Compared to other proposals for quantum computers, quantum dot-based
schemes have the important advantage of being compatible with conventional solid
state electronics. A severe problem, however, is posed by the relatively short decoherence times, which corrupt the desired evolution of the system states.
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Syntheses and Characterizations
3.1
Semiconductor Nanoparticles
3.1.1
Synthesis and Characterization of II–VI Nanoparticles
Alexander Eychmüller
The field of ‘‘synthesis’’ or ‘‘preparation’’ of II–VI semiconductor nanoparticles has
experienced an enormous development in the past two decades. The literature
dealing with this topic is virtually unmanageable. Nevertheless, the present chapter
attempts to give a chronological overview of some of the major historical lines in
this area, starting with the earliest work on CdS nanocrystals prepared in aqueous
solution. At several points in the story, mostly when successful preparations are
reported for the first time, the chapter branches into these evolving sub-fields and
leads to Sections 3.1.1.2 ‘‘Thiol-Stabilized Nanoparticles’’ and 3.1.1.3 ‘‘The ‘‘TOP/
TOPO’’ Synthesis’’. The remainder of this historical review then deals with matters
distinct from these preparational approaches. Sections 3.1.1.4 ‘‘Core-Shell Nanoparticles’’ and 3.1.1.5 ‘‘Quantum Dot Quantum Wells’’ deal with more complicated
nanoheterostructures in which two compounds are involved in the build-up of
spherically layered particles.
Naturally, in all chapters, the view of the author is indicated concerning the importance of work reviewed for the evolution of the subject as a whole. Besides this
necessarily subjective viewpoint, it is not impossible that other important work
may be overlooked. Thus, even if a particular chapter provides little more than a
list of literature references, only questions like ‘‘What has been done in this field’’
may be answered but never those like ‘‘What has not been done yet’’!
Historical Review
Even the shortest history of the preparation of II–VI semiconductor nanoparticles
has to start with the work published by A. Henglein in 1982 [1]. Like others [2–7],
this paper deals with surface chemistry, photodegradation, and catalytic processes
in colloidal semiconductor particles. But it is this very work which reveals the first
absorption spectrum of a colloidal solution of size-quantized CdS nanocrystals.
3.1.1.1
Nanoparticles: From Theory to Application. Edited by Günter Schmid
Copyright 8 2004 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
ISBN: 3-527-30507-6
3.1 Semiconductor Nanoparticles
Fig. 3-1
Absorption and emission spectra of a colloidal CdS solution [1].
These were prepared from Cd(ClO4 )2 and Na2 S on the surface of commercial silica
particles. As seen from Figure 3-1, the absorption onset is shifted considerably to
higher energies with respect to the bulk bandgap of CdS (515 nm). In addition, the
sol emits light upon excitation at 390 nm, which is also a matter of investigation in
this paper. The first correct interpretation of the observed blue shift of the absorption as a quantum mechanical effect is due to L. Brus [8]. In the framework of the
effective mass approximation, the shift in the kinetic energy of the charge carriers
due to their spatial restriction to the volume of the nanometer-sized semiconductors is calculated. (It is noted that comparable experimental [9] and theoretical
work [10] was carried out at almost the same time on the I–VII compound CuCl in
the Soviet Union.)
Another milestone in the preparation of II–VI semiconductor nanocrystals is the
work by Murray, Norris and Bawendi from 1993 [11]. Their synthesis is based on
the pyrolysis of organometallic reagents like dimethyl cadmium and trioctylphosphine selenide after injection into a hot coordinating solvent. This approach provides temporally discrete nucleation and permits a controlled growth of the nanocrystals. Figure 3-2 shows the evolution of the absorption spectra of a series of
CdSe crystallites ranging in size from about 1.2 to 11.5 nm. This series spans a
range of sizes from more or less molecular species to fragments of the bulk lattice
containing tens of thousands of atoms. The remarkable quality of these particles
is clearly seen in several features of the absorption spectra, e.g., the steepness of
the absorption onset, the narrow absorption bands, and the occurrence of higher
energy transitions.
By comparison of Figure 3-1 with Figure 3-2, it is evident that within about a
decade tremendous progress was made in the preparation of such systems. To illustrate what happened during this period of time, a few endeavors in both preparation and characterization from a number of different laboratories will be mentioned. Rossetti et al. reported on quantum size effects in the redox potentials and
electronic spectra of CdS nanocrystals [12]. Resonance Raman spectra were taken
from freshly prepared samples and from larger, ‘‘aged’’ particles. The size distributions were studied with transmission electron microscopy (TEM), and the crys-
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3 Syntheses and Characterizations
1993 state-of-the-art absorption spectra of CdSe
colloidal solutions (‘‘TOP-TOPO synthesis’’) [11].
Fig. 3-2
tallinity and crystal structure were determined. In a series of publications, Henglein and coworkers placed stress on the dissolution of CdS particles [13], enlarged
the field by the preparation of ZnS and ZnS/CdS co-colloids (including TEM characterization) [14, 15], gave insight into photophysical properties like time-resolved
fluorescence decays [16], and introduced the term ‘‘magic agglomeration numbers’’ [17]. Ramsden and Grätzel presented the synthesis of CdS particles formed
in aqueous solution from Cd(NO3 )2 and H2 S or Na2 S [18]. In combination with
luminescence data, transmission electron microscopic images were reported, and
the particles obtained were assigned to the cubic phase. In 1985, A. J. Nozik appeared on the scene with a paper on phosphate-stabilized CdS and small PbS colloids [19]. After yet another theoretical paper [20], it was again L. Brus who published in 1986 a review-like article [21] summarizing the then current status in the
field of semiconductor clusters, from both an experimental and a theoretical viewpoint. Also in 1986, attempts at a fractional separation by exclusion chromatography of an as-prepared sol of hexametaphosphate (HMP) (more correctly: Grahams
salt)-stabilized CdS particles were presented, together with stopped-flow experiments on extremely small clusters [22]. The binding of simple amines to the sur-
3.1 Semiconductor Nanoparticles
Absorption and luminescence spectra of CdS sols
(‘‘polyphosphate preparation’’) with varying starting pH [27].
Fig. 3-3
face of CdS particles of various sizes prepared by conventional routes as well as in
AOT/H2 O/heptane reverse micelles enhances their luminescence quantum yield
dramatically [23]. This effect is interpreted as a modification of mid-gap states
which may play an important role in the non-radiative recombination processes of
electron and hole in the materials under investigation. New preparative approaches
were presented, e.g., the preparation of CdS particles in dihexadecyl phosphate
(DHP) surfactant vesicles [24] and the synthesis of colloidal CdSe [25], which has
now become the most important model system for the investigation of the properties of nanocrystalline matter. As determined in pulse radiolysis experiments, an
excess electron on the surface of CdS colloids causes a bleaching of the optical absorption at wavelengths close to the onset of absorption [26]. One year later, a study
was published, likewise from the Henglein group, which since then has been cited
frequently [27]. CdS sols with particles ranging in size between 4 and 6 nm (as
seen from TEM-based particle size histograms) were prepared by first precipitating
Cd 2þ ions with the stoichiometric amount of injected H2 S followed by an ‘‘activation’’ of the colloid by addition of NaOH and excess Cd 2þ ions, resulting in a
strong emission band close to the energy of the band edge. Figure 3-3 reflects the
influence of the ‘‘starting pH’’ on the color of the emitted light after the ‘‘activation’’ process as well as on the particle size. The ‘‘activation’’ is interpreted in terms
of a precipitation of cadmium hydroxide onto the surface of the CdS particles. By
this procedure, the photostability of the sols is strongly increased approaching that
of organic dyes like Rhodamine 6G.
The reported fluorescence quantum yields, >50% in this publication and 80% in
earlier publications of other authors [23], seem over-optimistic. From our modern
viewpoint, the ‘‘activation’’ should rather be interpreted as being the result of the
creation of a shell of the large bandgap material Cd(OH)2 on the surface of the CdS
53
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3 Syntheses and Characterizations
particles (cf. also Section 3.1.1.4). This results in the formation of an electronic
barrier at the surface of the particles, preventing the charge carriers from escaping
from the particles’ core. This preparation operated as the basis for some novel coreshell structures and mulilayered nanotopologies, described in more detail in Sections 3.1.1.4 and 3.1.1.5.
An enormous step forward (especially in terms of the characterization involving,
among others, infrared and NMR spectroscopy, high-resolution electron microscopy (HRTEM), and X-ray diffraction (XRD)) was achieved by the work of the
AT&T group of Steigerwald and Brus from 1988 [28]. Organometallic compounds
were used to synthesize nanometer-sized clusters of CdSe in inverse micellar solutions. Chemical modification of the surface of these clusters was performed, the
cluster molecules could be removed and isolated from the reaction mixture, and
this was followed by a detailed powder XRD investigation. Later in the year, reports
on microheterogeneous systems appeared involving CdS, Znx Cd1x S, and CdS
coated with ZnS, which were formed in and stabilized by, again, DHP vesicles [29].
Additionally, layers of CdS particles (5 nm in size) were prepared by the LangmuirBlodgett technique [30]. These and many more results on small particle research
both on metal and on semiconductor particles are summarized in review articles
by Henglein [31, 32] and by Steigerwald and Brus [33]. Furthermore, considerable
progress has been made in the radiolytic production of CdS particles [34], in the
preparation of CdTe and ZnTe [35], and in the application of tunneling microscopy
to quantum-sized CdS [36]. The use of thiophenol as the stabilizing agent for CdS
clusters is reported by Herron et al. [37]. Competitive reaction chemistry between
CdS core cluster growth and surface capping by thiophenolate is assumed to be
responsible for the evolution of clusters whose core consists of essentially cubic
CdS and whose reactive surface has been passivated by covalently attached thiophenyl groups. The synthesis of porous quantum-size CdS membranes is achieved
in three subsequent steps: initial colloid preparation following standard procedures
[27], concentration of the suspension and dialysis, and finally the controlled evaporation of the residual solvent [38]. In a study of photoreductive reactions on small
ZnS particles, an elaborate HRTEM analysis was presented enabling the determination of the actual particle size by simply counting the lattice planes [39]. Postpreparative size separation of colloidal CdS by gel electrophoresis was demonstrated [40], and an ion-dilution technique was utilized in order to form ultrasmall
particles of CdS and CdSe in Nafion membranes [41]. After initial investigations
of the recombination processes of electron-hole pairs in clusters [16, 42–44], two
reports appeared on this topic, both of which proposed a mechanism of delayed
emission for the radiative ‘‘band-to-band’’ recombination [45, 46]. During this
time, several reports dealt with quenching and ‘‘antiquenching’’, i.e., the enhancement of the luminescence of nanosized CdS, all contributing in a qualitative manner to a deeper understanding of optical properties of suspended nanocrystals
[47–52]. CdS nanoparticles were synthesized in several non-aqueous solvents, and
their catalytic efficiency in the photoinitiation of polymerization of various vinylic
monomers were examined in detail [53].
Early in 1993, a special issue of The Israel Journal of Chemistry was devoted to
3.1 Semiconductor Nanoparticles
quantum-size particles and thus contained several articles which are concerned
predominantly with the synthesis and characterization of II–VI-semiconductor
nanoparticles [54–60]. Later in the year, two frequently cited articles summarized
the achievements in the field [61, 62]. The characterization of wet-chemically synthesized CdTe and HgTe as well as of ternary HgCdTe compounds (including XRD,
TEM and photoluminescence) were the focus of work by Müllenborn et al. [63].
The nature of the stabilizer plays an important role in the efficiency of luminescence quenching as demonstrated by the comparison of inverse micelle/HMPand thiophenol-capped CdS [64]. Thiol capping is also the clue to synthesizing
differently-sized CdTe particles prepared in aqueous solution [65]. In this interesting paper, the bandgaps of 2.0 nm, 2.5 nm, 3.5 nm, and 4.0 nm crystallites were
estimated from pulse radiolysis experiments, and their dependence on the size of
the particles was found to be in better agreement with pseudopotential [66] and
tight-binding models [67, 68] than with the effective mass approximation [20]. An
important result of the work of Rajh et al. is the observation that the extinction
coefficient per particle is independent of the size as predicted by theory [69–71].
This again brings us to the time of the work by Murray et al. [11]. Yet another aspect of its superior quality should now be mentioned, namely the thoroughness of
the characterization of the particles prepared. Among other findings, HRTEM images revealed the presence of stacking faults in certain crystallographic directions,
and showed that the particles are prolate and that the crystallites arrange themselves freely into secondary structures. The investigations probing only a limited
number of particles were supplemented by a carefully conducted X-ray diffraction
study including structural simulations. This work forms the basis of the ‘‘organometallic preparative route’’. Their most recent accomplishments are outlined in
Section 3.1.1.3, and some of the work on core-shell nanoparticles is described in
Section 3.1.1.4.
Like this approach, thiol stabilization proved to offer a very successful route
toward monosized nanoparticles of various compounds. Since this preparative
approach has also gained widespread interest, it will be outlined separately in the
next chapter.
From this point onwards in this ‘‘historical review’’ we will focus only on some
recent developments which differ from the ‘‘organometallic route’’ and from the
thiol technique.
Narrowing of the size distribution of a quite polydisperse CdS colloid by sizeselective photocorrosion was demonstrated by Matsumoto et al. [72]. In order to
replace the hazardous metal alkyls in the organometallic synthesis, a single-source
approach to the synthesis of II–VI nanocrystallites was presented by Trindade and
O’Brian [73]. CdSe particles were formed from methyldiethyldiselenocarbamato
cadmium (II), a compound which was synthesized beforehand from dimethyl
cadmium and bisdiethyldiselenocarbamato cadmium (II). Vesicles were again used
as reaction compartments [74], and nanophase reactors were generated by binary
liquids at the surface of colloidal silica particles also acting as reaction compartments for small CdS and ZnS particles [75]. The fabrication of composites from
CdSe nanoclusters and phosphine-functionalized block copolymers was the aim of
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3 Syntheses and Characterizations
Fogg et al. [76]. CdS particles formed from Na2 S and CdSO4 have been studied
extensively by 113 Cd and 1 H NMR [77]. From a careful analysis of the temperaturedependent NMR spectra, three different Cd 2þ species (located in the core, at the
surface, and in solution) are distinguished, and a considerable amount of water
has been found to be present in the dry samples. Several procedures are reported
to achieve a protection of the surface (by e.g. SiO2 [78]) or an encapsulation
[79, 80]. However, it is a complex interplay between the nature of the solvent, the
anions of the cadmium salt, the ratio of salt to thiurea, and the temperature which
controls the morphology and sizes of nanocrystalline CdS products. As shown by
XRD and TEM, spheres of various sizes, rods and even nanowires of up to 900 nm
in length were generated by a solvothermal process [81]. The luminescence properties of CdS particles prepared either in the presence of a poly(aminoamine) dendrimer or polyphosphate were revisited with special emphasis on the polarization
of the emitted light [82] and on the luminescence lifetimes [83]. In another report,
the charging of individual CdSe particles was studied [84]. Further synthetical efforts yielded HgE (E ¼ S, Se, Te) nanoparticles [85], PbS sols with various capping
agents like poly(vinyl-alcohol), poly(vinyl-pyrrolidone), gelatin, DNA, polystyrene,
and poly(methyl-methacrylate) [86], and ZnS nanoparticles in a silica matrix [87].
Small CdTe particles are accessible by a preparation in reverse micelles [88], and
‘‘peanut-like’’ nanostructures of CdS and ZnS were obtained by reactions at a static
organic/aqueous biphase boundary [89]. In 2001, the group of M. Lahav reported
on the organization of CdS and PbS nanoparticles in hybrid organic/inorganic
Langmuir-Blodgett films and the characterization of the layers by XRD- and TEMtechniques [90]. The zeolite MCM-41 is suitable as a host for CdS nanocrystals
which are formed via an ion-exchange reaction as shown by Zhang and coworkers
[91]. A completely new approach to the synthesis of Cd-chalcogenide nanoparticles
is demonstrated by Peng and Peng using CdO as the cadmium precursor and
strong ligands such as hexylphosphonic acid or tetradecylphosphonic acid [92].
Very fine (‘‘high-quality’’) nanocrystals are gained by this method, and even shape
control (i.e. spheres or rods) is demonstrated. Four types of binding sites at
the surface of CdS particles prepared by precipitation from a solution of cadmium carboxylate in dimethyl sulfoxide (DMSO) were identified by means of IR
and NMR spectroscopy, namely oxidized sulfur, adsorbed acetate, oxygen-bound
DMSO, and hydroxyl ions [93]. Micelles of various compositions are involved in
preparations of CdS (block-copolymer (compound) micelles [94, 95]) and triangularly shaped CdS nanocrystals (Cd(AOT)2 /isooctane/water reverse micelles [96]).
The protection of CdS particles by capping with a shell of silica was reported by
Farmer and Patten [97]. Furthermore, organic dendrons have been used for the
stabilization of CdSe quantum dots [98]. The optical properties of CdS quantum
dots in Nafion are the focus of a study by Nandakumar et al. [99], namely the twophoton absorption and the photon-number squeezing with CdSe particles in
PMMA [100]. Molecular species like (Li)4 -[Cd10 Se4 (SPh)16 ] as fragments of the
bulk structure act as templates for the growth of nanocrystalline CdSe in hexadecylamine as the solvent [101]. This approach may also be applied to the synthesis
of ZnSe and CdSe/ZnS core-shell nanocrystals.
3.1 Semiconductor Nanoparticles
Many of these efforts have also been included in review-like articles [102–108].
Empiricism, chemical imagination, and intuition largely form the basis for the
preparation of nanosized objects in solution. In the opinion of the author, this will
continue for a number of years despite attempts to successfully predict superior
preparative routes by means of combinatorial and computer scientific approaches
[109].
Thiol-Stabilized Nanoparticles
The use of thiols as capping agents in the synthesis of II–VI semiconductor nanocrystals has been mentioned several times in the preceding section. After some
preliminary work [110, 111, 54] based on earlier work by Dance et al. [112–114]
and others [115, 116], this approach came to full fruition for the first time in a
paper by Herron et al. [117]. This reported on the synthesis and optical properties of
a cluster with a 1.5 nm CdS core having the total formula [Cd32 S14 (SPh)36 ] 4 DMF
(SPh ¼ thiophenol, DMF ¼ N,N-dimethylformamide). Its structure, being a piece
of the cubic lattice as determined by single crystal XRD, will be discussed in
more detail in Section 4.2.1. In 1994 a paper was published on the synthesis,
the characterization, and several optical properties of thiol- and polyphosphatestabilized CdS particles ranging in size from 1.3 nm to about 10 nm as determined
from small- and wide-angle XRD as well as from TEM [118]. Furthermore, this
paper yields chemical and thermogravimetric analysis data, explains (also for CdS
small particles) the independence of the extinction coefficient per particle on the
particle size, plots the temperature dependence of the first electronic transition for
the differently sized particles, and reports the observation of a ‘‘reversible absorbance shift’’. This shift is detected as the difference between the transition energy
recorded in solution and that of a film formed from the particles on a quartz slide.
Initially, this was taken as an indication of a cluster-cluster interaction operative
in the films. One year later, Nosaka et al. reported on the formation process of
‘‘ultrasmall particles’’ from cadmium complexes of 2-mercaptoethanol [119]. In
this very thorough study, the stopped-flow technique was utilized to gain insight
into the very first steps of the particle evolution, and NMR yielded information
as to the presence of several different cluster species in solution amongst which
the [Cd17 S4 (SR)24 ] 2þ (RS ¼ 2-mercaptoethanol) exhibits the same tetrahedral core
structure as the crystallized [Cd17 S4 (SR)26 ] cluster of Vossmeyer et al. [120]. (cf.
Section 4.2.1).
The same preparative approach, i.e. the dissolution of a metal salt in water in the
presence of a stabilizing thiol, yields nanocrystals of CdSe [121], CdTe [122–124],
HgTe [125], and CdHgTe [126–128]. The thiols used were 1-thioglycerol, 2-mercaptoethanol, 1-mercapto-2-propanol, 1,2-dimercapto-3-propanol, thioglycolic acid,
thiolactic acid, and cysteamine [129]. Subsequent to the first steps in the preparation, heating of the reaction solution may be applied in order to initiate particle
growth. Thus, this preparative approach relies on the separation of nucleation and
growth in a way very similar to the organometallic TOP/TOPO (tri-octyl phosphine/tri-octyl phosphine oxide) preparation outlined in more detail in the following chapter [11]. Yet another similarity between the two strategies is the post3.1.1.2
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3 Syntheses and Characterizations
Powder X-ray diffractogram of HgTe nanocrystals
(‘‘thiol synthesis’’) with corresponding bulk values given as a
line spectrum [125].
Fig. 3-4
preparative treatment by applying size-selective precipitation to the crude cluster
solutions [130]. The sizes of the evolving nanocrystals depend mainly on the concentration ratios of the chemicals involved and on the duration of the heat treatment.
As examples of standard characterizations, some findings of studies on nanosized HgTe particles are shown [125]. Figure 3-4 shows the powder X-ray diffraction pattern (P-XRD) of a fraction of HgTe nanoparticles. The crystalline structure
derived from the positions of the wide-angle diffraction peaks is clearly the cubic
(coloradoite) HgTe phase. The wide-angle diffraction peaks are broadened because
of the small particle size. The mean cluster sizes estimated from the full width
at half-maximum intensity of the (111) reflection according to the Scherrer equation gives a value of about 3.5 nm. The weakly resolved small-angle X-ray diffraction peak indicates a rather broad size distribution of the HgTe nanocrystals
in the aqueous colloidal solution. In order to determine the size, shape, and crystal
structure of the particles, HRTEM images are taken. Figure 3-5 shows a typical
overview image of the sample together with a micrograph of an individual HgTe
particle with the corresponding Fast Fourier Transform (FFT). Like the X-ray diffraction data, the HRTEM images indicate a rather broad size distribution of the
HgTe nanoparticles, with sizes ranging from 3 to 6 nm. The existence of the lattice
planes on the HRTEM images of the particles indicates a high sample crystallinity. The images of single particles exhibit an interplanar distance of 3.73 Å, belonging to the (111) lattice plane of the coloradoite HgTe phase, which again is
consistent with the X-ray diffraction data. The photophysical properties of such
thiol-stabilized particles, such as photoluminescence [131] and optically detected
magnetic resonance (ODMR) [132], their dispersion in polymer films in order to
fabricate light-emitting devices [133–135], their possible use in telecommunication
3.1 Semiconductor Nanoparticles
High-resolution transmission electron microscopy of
thiol-stabilized HgTe nanocrystals; single particle in the inset
with 1st fast Fourier transform [125].
Fig. 3-5
devices [136], further structural properties unraveled by EXAFS (extended X-ray
absorption fine structure) [137–141], and their incorporation into photonic crystals
[142–144] and silica spheres [145] have been reported and have been in part already reviewed in several survey papers [146–149].
As an example of the capabilities of EXAFS spectroscopy, in Figure 3-6 the mean
CdaS distances as a function of the diameters of a series of CdS nanocrystals is
depicted [139]. These data are gained from a thorough temperature-dependent
study of the size dependence of various structural and dynamic properties of CdS
nanoparticles ranging in size from 1.3 to 12.0 nm. The properties studied include
the static and the dynamic mean-square relative displacement, the asymmetry of
the interatomic CdaS pair potential, with conclusions drawn as to the crystal
structure of the nanoparticles, the Debye temperatures, and the CdaS bond
lengths. As seen from Figure 3-6, the thiol-stabilized particles (samples 1–7) show
an expansion of the mean CdaS distance, whereas the phosphate-stabilized particles (samples 8–10) are slightly contracted with respect to the CdS bulk values.
The difference between the two series of nanocrystals consists in the type of interaction between the ligands and the Cd atoms at the surfaces. The thiols are covalently bound, and through this they have an influence on the structure of the
particles as a whole (most probably including steric requirements of the ligands
at the surfaces), whereas the phosphates are only loosely bound, which allows
the particles to contract slightly like a liquid droplet which tries to minimize its
surface.
59
60
3 Syntheses and Characterizations
Mean CdaS distances of thiol- and phosphatestabilized CdS nanocrystals determined by EXAFS at 5 K as a
function of the particle diameter [137–139].
Fig. 3-6
A large percentage of the atoms of particles in the nanometer size regime are
located at their surface, which makes investigations into the surface structure
highly desirable. The surface structure of thiophenol-stabilized CdS nanocrystals is
the subject of two articles (1997 and 1998) [150, 151]. Both research groups use
NMR techniques in order to study the functionalization, structure and dynamics of
the terminating ligands. Also, PbS [152] and ZnS [153, 154] nanocrystals may be
prepared with thiols as capping agents, showing again the versatility of this
approach. The interaction of g-radiation with Cd 2þ - and thiol-containing sols yields
small CdS particles, with the dose of the radiation determining the size of the
evolving particles [155]. The reverse process, i.e. the size-selective dissolution of
preformed CdS nanocrystals by photoetching, is demonstrated in a recent paper by
Torimoto et al. [156]. In an impressive study, Aldana, Wang and Peng give detailed
insight into the photochemical degradation of thiol-stabilized CdSe particles [157].
At least three steps are involved in this process, namely photocatalytic oxidation of
the thiol ligands on the surface, photooxidation of the nanocrystals, and finally
precipitation of the nanocrystals. A bichromophoric nanoparticle system consisting
of thioglycolic acid-stabilized CdTe particles and layer-by-layer assembled films of
an anionic polyelectrolyte is presented by Westenhoff and Kotov [158]. The first
investigation into the electronic structure of a nanoparticle dimer was reported by
Nosaka and Tanaka [159]. The semiempirical molecular orbital calculations first
investigated the influence of phenyl-capping of polynuclear cadmium complexes
and subsequently yielded negligible influence on the excited state by dimer formation.
In a recent report, Gaponik et al. stress the advantages of thiol-capping of CdTe
as an alternative to organometallic preparation routes [160]. This article reports on
3.1 Semiconductor Nanoparticles
the aqueous synthesis, the structure, some optical properties, and the processability of thiol-capped CdTe nanocrystals, with an emphasis on the procedures,
leading to highly luminescent particles having room temperature quantum efficiencies of up to 40%. The advantages and disadvantages of different thiols used as
capping agents of CdTe QDs which are relevant for the further design of surface
properties of the nanocrystals are discussed. Summarizing the experimental data
available today, thiol-capped CdTe nanocrystals synthesized in aqueous solution
represent a kind of highly luminescent core-shell QD system with a naturally
sulfur-capped surface (CdS shell) created by mercapto-groups covalently attached to
the surface cadmium atoms. In yet another article, the stability of thiol-stabilized
semiconductor particles is investigated [161]. Various analytical methods like Analytical Ultracentrifugation, UV-vis absorption spectroscopy, NMR spectroscopy, and
powder X-ray diffraction are applied. Reduced particle sizes are found in solution
at low particle concentrations. Most probably, the assumed covalently bound thiols
desorb from the surface of the particles, leaving behind vulnerable unstabilized
particles that might even undergo continuous decay over time. Additionally,
breaking of the intra-ligand SaC bond is observed, resulting in a cadmium sulfide
particle synthesis introducing sulfur only by the ligands, without any additional
sulfide ions.
Finally, it may be noted that the nanocrystals prepared in aqueous solutions may
be transferred into different nonpolar organic solvents [162]. This is accomplished
by a partial ligand exchange, with acetone playing an important role in the efficient
phase transfer. Since the nanocrystals retain their luminescence properties after
being transferred, this procedure provides a new source of easily processable luminescent materials for possible applications.
The ‘‘TOP/TOPO’’ Synthesis
The term ‘‘TOP/TOPO’’ synthesis refers to an organometallic approach to synthesizing CdE (E ¼ S, Se, Te) nanocrystals first introduced in 1993 by Murray et al.
[11]. The preparative route is based on the pyrolysis of organometallic reagents
(like dimethylcadmium and bis(trimethylsilyl)selenium) by injection into hot coordinating solvents (like tri-n-octylphosphine oxide (TOPO) and tri-n-octylphosphine
(TOP)). This provides temporally discrete nucleation and permits the controlled
growth of nanocrystals. Subsequent size-selective precipitation isolates samples
with very narrow size distributions. The widespread success of this preparation is
largely due to its versatility, its reproducibility and, last but not least, the high
quality of the particles in terms of crystallinity and uniformity. The versatility is
seen, e.g., in the successful transformation of the preparation principles beyond
the II–VI materials to the classes of the III–V and IV–VI semiconductor nanocrystals [163–172]. It is probably not an exaggeration to identify the first reports on
the organometallic synthesis also as the trigger for elaborate theoretical work on
size-dependent properties of semiconductor clusters like, e.g., the electronic structure (see, e.g., [173–183]). The years immediately after the ‘‘invention’’ of the TOP/
TOPO preparation are characterized by investigations into the optical properties of
CdSe particles of a wide range of sizes and the surface properties of the evolving
3.1.1.3
61
62
3 Syntheses and Characterizations
nanocrystals [184–195]. Most of this beautiful work has been presented in a very
readable review by the founders of the field [196], including the outstanding work
on Förster energy transfer taking place in close packed assemblies of CdSe quantum dot solids [197, 198]. The photostability and the high luminescence quantum
yield made the particles ideal candidates for single-particle emission studies [199–
212].
Also, some of the early work on ‘‘high-tech’’ applications of semiconductor
nanoparticles is based on the above-mentioned preparative route, namely the construction of a light-emitting diode made from CdSe nanocrystals and a semiconducting polymer [213–215], and even a single electron transistor is made by the
use of CdSe nanocrystals [216].
The charge-tunable optical properties of CdSe particles derived by the TOP/
TOPO method are studied [217], and these are used as precursors for the generation of thin films of CdTe [218]. Other effects like the photoionization in CdTe
particles embedded in TOPO are investigated, elucidating some of the photophysical observations made in single-particle studies [219]. N-type colloidal semiconductor particles of CdSe are made via an electron transfer process [220], and
the synthesis of ZnSe nanocrystals [221] and their doping with manganese is reported [222]. The fate of the photogenerated charge carriers is the topic of several
studies of particles derived by the organometallic procedure (see, e.g., the review
[223] and references therein and [224–231, 100]), including the size-dependent
photoconductivity of CdSe nanoparticles [232]. Spin spectroscopy of ‘‘dark excitons’’ (see also the contribution of U. Banin in this book) in CdSe particles of
various sizes is performed as a function of temperature and magnetic field
strength [233]. The influence on the sizes, size distributions, and luminescence
quantum yields of various amines on addition to the reaction mixtures is studied
by Talapin and coworkers [234–236]. Alternative routes toward high quality CdSe
nanocrystals are discovered by Qu, Peng and Peng by varying the solvents/ligands
and the precursors like CdO and Cd(Ac)2 instead of Cd(CH3 )2 [92, 237].
Organometallically synthesized CdSe (and InAs) nanocrystals act as experimental examples in a very interesting study concerning the kinetics of nanocrystal
growth [238]. It is demonstrated that diffusion-limited growth can lead to narrowing of size distributions with time and that, under certain conditions, the kinetics
of nanocrystal growth may also be influenced by the variation of the surface energy
with size. (The matter of kinetic-dependent crystal growth is also considered for
size-tunable CdS particles prepared by other routes [239] and for nanometer sized
crystalline particles in glassy media [240].) In a theoretical study, the ideas of Alivisatos [238] are reconsidered, giving even more detailed insight into the evolution of ensembles of nanoparticles in colloidal solutions [241]. An outline is given
of how the growth of nanoparticles in the diffusion-controlled regime may result
in better final size distributions than those found for nanoparticles grown in the
reaction-controlled regime. Additionally, possible ways of controlling the particle
size distributions in colloidal solutions are discussed. These include, for example,
the increase in the surface tension at the solvent-nanoparticle interface and the
decrease in the diffusion or mass transfer coefficient. A related experimental study
3.1 Semiconductor Nanoparticles
is published by Qu and Peng in 2002 [242] in which the photoluminescence quantum yield (PL QY) of organometallically prepared CdSe nanocrystals is studied
during their growth. A ‘‘bright point’’ is observed which denotes the fact that under
a given set of initial conditions of the colloidal solution the PL QY increases to a
maximum value before decreasing again. Very similar results are reported in an
even more thorough investigation by Talapin et al. [243]. Results concerning the
dynamic distribution of growth rates within ensembles of CdSe and InAs nanoparticles prepared organometallically and CdTe prepared in aqueous solutions are
presented indicating that the observed phenomenon of fraction-dependent PL QY
is of more general character. Relating the results to the above-mentioned theoretical work, it is proposed that the nanocrystal fractions with the best surface quality
(in terms of a minimum of defects) correspond to particles with the smallest net
growth rate within the ensemble because annealing of surface defects is most effective at this ‘‘zero growth rate’’. In a further study stretching beyond the topic of
this chapter, X. Peng et al. report even on shape control over CdSe nanoobjects by
proper adjustment of the initial experimental conditions [244] (see also [245, 246]).
In two recent studies, the absorption cross section (or extinction coefficient) of
CdSe nanocrystals has been under investigation [247, 248]. Striolo et al. [247] report (among other results) that the oscillator strength of the first exciton transition
per CdSe unit is more or less independent of particle size, in accord with Schmelz
et al. [249]. In contrast, Bawendi and coworkers clearly show that as the quantum
dot radius decreases, the oscillator strength per CdSe unit of the lowest-energy
transition increases (as reported by others for CdS [118] and CdTe [65] prepared
differently), whereas the spectra converge for higher energies. Thus, this issue remains to be elucidated further.
Core-Shell Nanocrystals
Coating a given ensemble of nanoparticles by another material yields ‘‘core-shell’’
nanocrystals. In these structures, the cores may be any kind of colloidal particles,
i.e. metals, insulators, and all classes of semiconductors. Likewise, the shells may
consist of any sort of materials including organic ones. This survey concentrates on
semiconductor cores and inorganic shells mainly of semiconductor materials.
In the majority of cases, the electronic structure of the core shell particles is as
follows. The core material having a certain bandgap is capped by a material with a
larger bandgap, the conduction band energy of the capping material being higher
(more negative in an electrochemical sense) than that of the core material and the
valence band energy of the capping material being lower (more positive on the
electrochemical energy scale) than that of the core material. This energetic situation is frequently called a type-I structure. The main consequence of this capping
is that the exciton photogenerated in the core is prevented from spreading over the
entire particle, and by this it is forced to recombine while spatially confined to the
core. In most cases this is accompanied by enhanced luminescence.
Early reports on the formation of coated nanoparticles include the aqueous
polyphosphate-stabilized CdS/Cd(OH)2 system [27]. The preparative procedure is
to add to a given sol of CdS particles a large excess of Cd ions, while increasing the
3.1.1.4
63
64
3 Syntheses and Characterizations
pH to achieve strongly basic conditions. The obtained increase in luminescence
quantum yield in comparison with the uncoated CdS particles is explained by saturation of free valences at the particle surface, and this is taken as an indication of
the formation of the proposed structure. The report is speculative in a sense that a
structural characterization is not performed. This is not very surprising, since obtaining clear-cut evidence for shell formation of a few monolayers of material on
nanoobjects is a characterizational challenge. This challenge has been accepted
more or less since then in a wealth of reports on systems like Ag2 S on CdS [250],
ZnS on CdS [29], Ag2 S on AgI [251], CdSe on ZnS [252], CdSe on ZnSe [253],
HgS on CdS [254], PbS on CdS [255], CdS on HgS [256], ZnSe on CdSe [257], ZnS
on CdSe [258–261], CdS on CdSe [262] and vice versa [263]. Many of these endeavors have already been reviewed [32, 62, 61, 103, 102, 147, 264–267] and will
therefore not be mentioned in detail. A very fine example of the most careful
characterization of the overgrowth of CdSe nanocrystals with ZnS is reported in
the paper of Dabbousi et al. (1997) [259]. For the shell growth on the preformed
CdSe nanocrystals, diethylzinc and hexamethyldisilathiane are used as Zn and S
precursors. Dissolution of equimolar amounts of the precursors in suitable solvents is followed by dropwise addition to vigorously stirred CdSe colloidal solutions
held at temperatures between 140 and 220 C, depending on the core sizes. Besides optical spectroscopy, the authors managed to perform wavelength dispersive
X-ray spectroscopy, X-ray photoelectron spectroscopy, small- and wide-angle X-ray
scattering, and (high-resolution) electron microscopy. As an example of the thorough characterization, Figure 3-7 is taken from this work. It demonstrates the Xray scattering in the small-angle region of bare 4.2 nm CdSe dots (curve a) and the
increase of the dot diameters with the increase in the ZnS shell thickness from
0.65 monolayers (curve b) to 5.3 monolayers (curve e). The gradual shift of the
ringing pattern to smaller angles indicates that ZnS is in fact growing onto the
CdSe dots instead of nucleating separate ZnS nanocrystals. This is also deduced
from the size histograms gained from electron microscopy of the same series of
particles. In Figure 3-8, the uniform growth of the particles is clearly demonstrated, together with a broadening of the size distribution. The aspect ratio histograms (right hand side of Figure 3-8) of the slightly elongated particles are monitored as well and are used to generate accurate fits to the small angle X-ray data
shown above. In an effort to relate the structural to the optical properties, the
growth of the ZnS shell on the CdSe cores is described as locally epitaxial. The ZnS
shell tends to create defects at coverages of more than 1.3 monolayers, accompanied by a decrease in luminescence quantum yield.
Based on this work, progress has been made with similar particles, e.g., in understanding the photo- and electrodarkening effects [260, 268] and in the assembly
of luminescing semiconductor nanocrystals with biological objects in order to
create fluorescent probes in sensing, imaging or other diagnostic applications
[269–271]. Just to name a few, in other reports using the described synthesis of
CdSe/ZnS core-shell particles, high-resolution spectral hole burning is performed
[272] or effects like photooxidation and photobleaching are investigated on the
single particle level [273]. Likewise, the emission properties of single ZnS or CdS
3.1 Semiconductor Nanoparticles
Small-angle X-ray scattering data from bare 4.2 nm
CdSe particles (a) overgrown by 0.65 (b) to 5.3 (e) monolayers
of ZnS [259].
Fig. 3-7
overcoated CdSe particles are investigated by several groups [200, 202–204, 206,
207, 209–211] (summaries in e.g. [212, 147]).
Other reports include structures in which either the core or the shell is a II–VI
semiconductor material such as CdS capped by sulfur or polysulfide [274, 275],
InAs capped by CdSe [276], ZnSe and ZnS [277, 278], or InP capped by ZnCdSe 2
[279] or ZnS (see, e.g., [280]).
New synthetic approaches include the formation of core-shell type particles from
CdTe and CdS [281], ZnSe on CdSe [282, 283], and CdS on HgTe [284]. In the
latter report, a synthetic route to strongly IR-emitting nanocrystals is outlined
whereby in this case capping enhances the robustness of the particles rather than
the (already high) luminescence quantum yield.
Quantum Dot Quantum Wells
As an extension of the work of precipitating HgS on polyphosphate-stabilized CdS
and CdS on HgS [254, 256], a wet chemical synthesis has been developed which
yields a nanoheterostructure consisting of three domains, viz., the system CdS/
HgS/CdS. The basic preparative idea involves the substitution reaction
3.1.1.5
ðCdSÞm þ nHg 2þ ! ðCdSÞmn ðHgSÞn þ nCd 2þ
(1)
65
66
3 Syntheses and Characterizations
Size histograms of a series of CdSe nanocrystals
overgrown with 0.65 (b) to 5.3 (e) monolayers of ZnS (A) and
corresponding aspect ratios (B) [259].
Fig. 3-8
taking place at the surface of the preformed CdS particles. This reaction proceeds
because of the at least 22 orders of magnitude smaller solubility product of HgS
compared to that of CdS. The art consists in managing the reaction to take place
only at the surface, which is largely achieved by strictly controlling the experimental conditions regarding pH, concentrations, duration of the addition, stirring, etc.
After a monolayer of HgS is built on the CdS cores, this structure is capped by
more CdS via the addition of H2 S and further Cd ions. The first reports deal with
structures having a CdS core of about 5 nm surrounded by a shell of 1–3 monolayers of HgS, which again is coated by up to five monolayers of CdS [285, 286].
b-HgS with a bulk band gap of 0.5 eV forms a quantum well inside the cubic CdS
(bulk band gap 2.5 eV), which forced the authors to describe these new structures
as ‘‘quantum dot quantum wells’’ or briefly QDQWs. Probably, an even better term
would be ‘‘quantum well in a quantum dot’’. Colloquially, the term ‘‘nano-onions’’
has been established meanwhile. The first four papers on these particles were
published in 1993 and 1994 and describe the synthesis, the characterization [285,
286], transient bleaching after laser excitation [287], and a widely applied theory
3.1 Semiconductor Nanoparticles
Models (a1, b, c1, and d1), HRTEM
micrographs (a2, c2, d2) and simulations
(a3, d3, and d4) of CdS and CdS/HgS nanoheterostructures in the course of preparation.
Fig. 3-9
The CdS surface is replaced by HgS (a1 to b).
Covering this structure with CdS leads to
epitaxial growth (b to c1) or to twinned
epitaxial growth (b to d1) [289].
based on the ‘‘effective mass approximation’’ [288]. As with the other findings, also
the bleaching experiments on the nanoheterostructures, make them to seem to be
homogeneous electronic systems, and the results are not explainable by the superposition of the properties of the segments. Together, these efforts are summarized
in review articles [103, 289, 290]. Reference [289] also contains the results of Mews
and Alivisatos [291, 292], including an impressive HRTEM investigation (see also
[105]) with accompanying simulations presenting information about the shape and
the crystallinity of the particles (cf. Figure 3-9). Hole burning and fluorescence line
narrowing experiments provide evidence on the degree of charge carrier localization in the HgS layer. Theoretical papers [293, 294] (likewise described in [289])
refer to these experiments and deduce that a ‘‘dark exciton’’ represents the lowest
excitation state in these nanoheterostructures, too. A recent tight-binding study by
Perez-Conde and Bhattacharjee also detects bright and dark states together with
the strong localization of electron and hole states within the HgS well [295]. The
tight-binding approach is also found by Bryant and Jaskolski to provide a good description for nanosystems with monolayer variations in composition [296]. It is
pointed out that trap states cannot be described solely as states in the quantum
well. Instead, depending on the shell thickness and band offset, trap states can
have a large interface character.
Two articles by El-Sayed and coworkers [297, 298] published in 1996 and 1998
prove experimentally that the electronic structure is more complicated and can be
67
68
3 Syntheses and Characterizations
observed by advanced spectroscopic techniques. The authors carry out experiments
analogous to the transient bleaching at CdS and CdS/HgS/CdS particles, but with
femtosecond time resolution. By this means it is possible to observe the kinetics of
the charge carrier capture in traps and in the HgS quantum well, respectively. Obviously, the spectral diffusion of the bleaching (of the ‘‘optical hole’’, as it is called
by the authors) takes place over an intermediate state to which the authors assign
an energetically relatively high ‘‘dark exciton’’ state predicted by theory [294]. In
this state, the electron is situated in the HgS well, whereas the hole is in the CdS
cladding layer of the particle and the ‘‘well’’ has a rather barrier-like effect on the
hole. In a subsequent study, similar experiments were carried out on particles with
a two-monolayer-thick well of HgS, giving good agreement with theoretical predictions regarding the exciton energies [299]. The dynamics of the localization of
the exciton in CdS/HgS QDQWs is also the topic of an investigation published by
Yeh et al. [300]. By comparing the results from nanosecond hole-burning experiments with those having femtosecond time resolution, the authors conclude that
the primary optical interaction excites electrons from a delocalized QDQW ground
state and not from a localized HgS well state. Only after longer times, of the order
of 400 fs, will the exciton finally be localized in the HgS well. ODMR (optically
detected magnetic resonance) is applied to QDQW samples with varying layer
thickness of HgS as well as of the outer CdS by Lifshitz and coworkers [301]. Because of its sufficient longevity, it is possible to study by this technique especially
the trap luminescence, which, compared to the near band edge fluorescence,
occurs red-shifted by about 400–700 meV. The change in luminescence which is
induced by magnetic resonance in the excited state is detected. From measurements carried out in dependence on the intensity of the luminescence excitation
and on the frequency of the coupled microwaves, with subsequent simulations of
the spectra, it may be concluded that there are at least two different recombination
channels via traps in the particles. These traps are localized in the CdSaHgS interfaces and could possibly be twin grain boundaries and edge dislocations. In a
photoluminescence experiment, an enhancement of the band edge emission was
observed upon excitation with two differently colored beams instead of one in an
ordinary experiment [302]. This result is explained as being the consequence of a
partial quenching of the defect emission, suggesting a mutual interaction of the
two luminescence events.
A strong reduction in the spectral and intensity fluctuations is observed when
QDQWs of CdS and HgS are compared with pure CdS particles in single-particle
emission studies [303]. This is, again, explained by charge carrier localization in
the HgS region of the nanocrystals.
A QDQW with a large lattice mismatch, namely ZnS/CdS/ZnS, was prepared by
the El-Sayed group [304]. Established techniques were combined to synthesize
small seeds of CdS and ZnS, and afterwards nanoepitaxy was applied to grow ZnS/
CdS nanoheterostructures. The evolving absorption spectra of the particles are
presented together with a theoretical treatment.
Very recently, the same group has published first results on a QDQW system
that contains two wells of HgS separated by a wall of CdS. The two-well system is
3.1 Semiconductor Nanoparticles
CdS
HgS
CdS
HgS
CdS
a b c d
Fig. 3-10 Schematic drawing of the double-well nanoheterostructures described in the text. The wells (a and c)
consist of the small bandgap material b-HgS. The core, the
barrier between the wells (b) and the cladding layer (d) are
made from CdS.
characterized by absorption and emission spectroscopy, clearly revealing the formation of a two-well system rather than a single double-layer structure [305].
In an article in the same year, Dorfs outlined efforts aimed in the same direction,
namely the preparation and characterization of a series of QDQWs containing two
wells [306]. The preparative route firstly follows the established procedures which
yield single-well particles consisting of a CdS core of diameter 4.7 nm surrounded
by one or two monolayers of HgS capped by two, three or four layers of CdS. Following the same concept as that for the single-well particles, i.e. repetition of the
replacement of the surface CdS by HgS, addition of a second monolayer of HgS
and capping of the whole structure by three monolayers of CdS yields a whole new
family of particles. As a nomenclature, the authors suggest for these nanoheterostructures the following (cf. Figure 3-10): CdS/HgS is represented by ‘‘abcd’’,
where, starting with the CdS core, ‘‘a’’ is the number of HgS monolayers, ‘‘b’’ is
the number of monolayers of CdS as the barrier between the wells, ‘‘c’’ is the
number of monolayers of the second HgS well, and ‘‘d’’ is the number of cladding
layers.
In Figure 3-11 are summarized the absorption spectra of a total of eight different
double-well QDQWs prepared according to the above outlined procedures. All
eight samples have in common the core size 4.7 nm and the outer cladding of
three monolayers of CdS. They differ in the three inner compartments: the inner
well consists of either one or two monolayers of HgS, the barrier between the wells
are either two or three monolayers wide, and the outer well again consists of one or
two monolayers of HgS. Thus, the structures CdS/HgS 1213, 1223, 1313, 1323,
2213, 2223, 2313, and 2323 are generated. The most remarkable findings gained
from the absorption spectra and the accompanying TEM study are as follows:
69
3 Syntheses and Characterizations
Absorption spectra of eight different double-well
nanoheterostructures. In CdS nanocrystals two wells of HgS are
embedded. Well thickness and well separations vary. For
nomenclature see text [306].
Fig. 3-11
1. Judging from the particle size histograms, there is no indication of separately
formed CdS or HgS. Exclusively homogeneous particle growth is observed.
2. The position of the first electronic transition is further into the red spectral
region the more HgS is present in the structures.
3. Comparing the four pairs of samples having the same well composition, the
transition is located further into the red when the wells are separated by only
two monolayers of CdS instead of three, which is explained by an enhanced
electronic coupling on decreasing the well separation.
4. The onset of absorption of the particles exhibiting an asymmetric inner potential (e.g. CdS/HgS 2313 and 1323) takes place at the same energy.
5. The high-energy absorption depends on the total amount of material present in
solution.
Some of these findings have already been treated theoretically, giving insight
into very interesting electronic properties and charge separation features [307].
400
Cd 3d52
hn = 494.7 eV
kcps
70
300
v
s
409
408
407 406 405 404
binding energy [eV]
403
residual plot
200
High-resolution XPS spectrum of the Cd 3d5/2 core
level in the CdS/HgS/CdS/HgS/CdS double-quantum-well
nanocrystals at a surface-sensitive energy [308].
Fig. 3-12
3.1 Semiconductor Nanoparticles
In an investigation by photoelectron spectroscopy (XPS) with tunable synchrotron radiation, a member of this family of particles was compared to single-well
CdS/HgS/CdS QDQWs [308]. As is to be expected, the photoemission spectra of
the CdS/HgS/CdS/HgS/CdS double-quantum-well nanocrystals look quite similar
to those of the single-well species. As seen in Figure 3-12, in addition to a bulk
species, high-resolution spectra also reveal a surface environment for Cd, while
different S species could not be resolved. To some extent the XPS experiments allow the characterization and verification of the onion-like structure of the quantum dot quantum well nanocrystals, which is shown also by an extension of a previously developed model for core-shell nanocrystals [309].
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3.1.2
Synthesis and Characterization of III–V Semiconductor Nanoparticles
Uri Banin
Introduction
The preparation of nanocrystals of III–V semiconductors presented a significant
and important challenge to researchers in the field for over a decade [1–7]. Many of
the III–V semiconductors are direct gap semiconductors, unlike silicon. This has
led to their widespread implementation as the central building blocks of present
day optoelectronic devices such as diode lasers, ubiquitous in telecommunication
3.1.2.1
79
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3 Syntheses and Characterizations
applications, providing obvious technological motivation for development of wet
chemical synthesis approaches for preparing nanocrystals of these materials. Furthermore, the opportunity to develop such routes for preparing III–V semiconductor nanocrystals presents a challenge for synthetic chemistry. While II–VI
semiconductors can be prepared directly in aqueous media using ionic precursors,
no equivalent synthetic procedures could be developed for the III–V semiconductor compounds. This results from the more covalent bonding typical of III–V
compounds and the high energetic barriers in the formation of the semiconductor
from suitable precursors.
The availability of high quality III–V semiconductor nanocrystals has contributed significantly to fundamental studies of the size-dependent electronic, optical,
electrical, and structural properties of nanocrystals [8–10]. An interesting case is
that of InAs, a narrow-band semiconductor (bulk band gap Eg ¼ 0:42 eV). Under
the conditions of quantum confinement, the band gap of InAs nanocrystals can be
tuned over a broad range in the near infra-red, where the selection of traditional
organic dyes is poor and their stability is low. This points toward potential applications of such nanocrystals in optoelectronic devices relevant for telecommunications [11] and as biological fluorescent markers in the near IR spectral region [12].
As III–V semiconductor QDs (for example, InAs) can also be prepared via a
strain-induced growth mechanism using molecular beam epitaxy (MBE) [13, 14],
there is an opportunity for direct comparison of the properties of two QD systems
– colloidal grown nanocrystals versus MBE-grown dots [15].
Early attempts to prepare III–V semiconductor nanocrystals include the synthesis of GaAs nanocrystals in 1990 by Olshavsky et al. [16] using a dehalosilylation
reaction of GaCl3 and As(Si(CH3 )3 )3 developed by Wells [17], where quinoline was
used as both the solvent and as the surface passivating ligand. These reactions
produced nanocrystals about 3–4 nm in diameter, but the optical absorption of the
particles showed some interference due to suspected Ga-quinoline complexes or
quinoline oligomeric species [18]. To date, most work on the preparation of III–V
nanocrystalline materials has been based on similar dehalosilylation reactions,
with variation of reactants, reaction conditions, and solvents.
An additional route for the preparation of GaAs, was reported by Nozik et al.
[19], who used Ga(acac)3 and As(Si(CH3 )3 )3 with triglyme as the solvent to produce
GaAs nanocrystals. However, the absorption spectra did not show a well-defined
feature that shifts with particle diameter, indicating a poor size distribution or
other problems with nanocrystal quality. Wells et al. [20, 21] have also prepared
nanocrystalline GaAs as well as GaP using the metathetical reaction of GaCl3 and
(Na/K)3 E (E ¼ P, As) in glymes. These reactions produced nanocrystalline material, but there is no separation of sizes, and the optical spectra do not exhibit the
characteristic well-defined absorption feature.
InP nanocrystals were prepared by Nozik et al. using the dehalosilylation reaction of an uncharacterized InCl3 -oxalate complex and P(Si(CH3 )3 )3 with trioctylphosphine oxide (TOPO) serving as the solvent and surface-passivating group [4].
This synthesis procedure provides high quality nanocrystals, and a size-dependent
feature is observed in the absorption spectra. Nozik also reports the synthesis of
3.1 Semiconductor Nanoparticles
nanocrystals of the indirect gap semiconductor GaP and the ternary material
GaInP2 using analogous reactions. Additionally, Guzelian et al. reported the successful synthesis of size-controlled InP nanocrystals [3]. Here, InCl3 dissolved in
TOP reacts with P(Si(CH3 )3 )3 in TOPO. The preparation of InAs nanocrystals is an
example of a synthesis procedure where a very high level of control on particle size
and monodispersity has been demonstrated [5]. We therefore choose to focus on
this particular system to obtain a more detailed view of the reaction procedure.
Recently, Nozik and coworkers also reported the synthesis of colloidal GaN nanocrystals. They used a polymeric gallium imide precursor {Ga(NH)3=2 }n , reacted in
trioctylamine at 360 C to produce GaN nanocrystals [22].
Following the improvements in the synthesis of InAs and InP nanocrystals,
workers in the field have implemented several strategies to improve the fluorescence and stability of the functional particles. The optical and electronic properties
of the particles can be further controlled by fabricating advanced core/shell nanocrystal structures, where a shell of a higher band gap semiconductor is grown on
the core, yielding superior surface passivation, as demonstrated first for II–VI
semiconductor nanocrystals [23–26] and more recently for III–V nanocrystals.
Synthesis of core/shell nanocrystals with InAs cores was recently developed and
will be outlined here [24–26]. The preparation of InP/ZnS core/shells with improved luminescence properties was reported by Weller and coworkers [27], while
Nozik and coworkers reported the synthesis of lattice-matched shells of ZnCdSe 2
on InP cores [28]. Another approach to improving the luminescence of InP nanocrystals was reported by Nozik and coworkers using etching with HF [29].
The chapter begins with a general discussion of the synthetic strategy employed
for III–V semiconductor nanocrystals. This is followed by a description of the synthesis of InAs nanocrystals as a prototypical example, and the characterization of
the particles by various methods is briefly reviewed. We then detail the synthesis of
core/shell nanocrystals with III–V semiconductor cores, focusing particularly on
InAs cores.
Synthetic Strategy
The first step in the synthesis of a nanocrystal is the nucleation of the particles.
Successful nucleation is an indication of whether the reaction in question is appropriate for the formation of nanocrystals. For a monodisperse sample, the nucleation event must be well separated in time from the growth step, which generally means that nucleation must occur on a short time scale, of the order of a
fraction of a second. In contrast, if nucleation is spread out in time such that nucleation and growth overlap, different nucleation sites will undergo growth of different durations, resulting in a broad size distribution.
An effective way to achieve separation of nucleation from growth is the use of a
method of injecting suitable precursors into the solvent at high temperature such
that nucleation occurs immediately on contact with the hot solvent. This method
has been perfected in the preparation of CdSe nanocrystals [30, 31] and was also
implemented for InP and InAs nanocrystals [3–5]. Difficulties in fully realizing the
benefits of the injection method in the case of III–V semiconductor nanocrystals
3.1.2.2
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3 Syntheses and Characterizations
include the need for prolonged reaction times to overcome the high reaction barriers and the need for annealing at high temperature to achieve better crystallinity
of the nanocrystals.
The next stage in the nanocrystal synthesis is the growth of the particles. As with
nucleation, intermediate species and the growing particle must stay in solution
throughout the process. Growth is completed when the reagents are consumed,
although additional high-temperature annealing time may be needed to obtain
good crystallinity. Under prolonged high-temperature annealing, the possibility of
Ostwald ripening arises – an undesired mechanism that will broaden the size distribution [32]. Growth can also be effectively quenched by cooling the reaction
mixture. A typical behavior during the growth stage was found by Alivisatos and
coworkers, who studied the growth of both CdSe and InAs nanocrystals and introduced the important concept of size focusing [31]. This concept is based on the
existence of a critical size that depends on the reaction conditions (concentration of
reactants, temperature, solvent etc.). Above the critical size, growth rate decreases
with increased size, leading to a narrowing or focusing of the size distribution,
while below the critical size, defocusing results because the growth rate of smaller
particles is slower than that of the bigger particles. To achieve a narrow distribution, conditions should be optimized so that during growth the reaction remains
within size focusing conditions.
An additional critical issue for high quality nanocrystals is the passivation of the
surface. During the synthesis, surface passivation helps to control the growth of
the nanocrystal, prevents agglomeration and fusing of particles, and provides for
the solubility of the nanocrystals in common solvents. Solvation is important not
only for the nanocrystals, but also for the reactants and intermediates. Surface
passivation is achieved by using an organic molecule to coordinate or bond to the
nanocrystal surface. Most commonly, this organic molecule is the solvent for the
reaction. This strategy benefits from having the capping molecule in huge excess
and from having a common solvation medium for the reactants, intermediates,
and nanocrystals.
Capping groups typically contain an electron-rich donating group such as phosphine oxides, phosphines, amines, or thiols and behave as a Lewis base, which coordinates to the electron-poor Lewis acid-like metal of the semiconductor such as
cadmium or indium. This coordination passivates the dangling orbitals at the
nanocrystal surface, preventing further growth or agglomeration. The other end of
the ligand imparts solubility to the nanocrystal by giving the particle a hydrophilic
or hydrophobic surface. For example, the alkyl groups of TOPO or TOP result in
nanocrystals that are soluble in relatively nonpolar solvents, while using ligands
such as mercaptobenzoic acid renders the particles soluble in polar solvents like
methanol [33].
Surface passivation plays an important role not only during the synthesis. It can
also be used to further derivatize the surface of the particles, providing additional
chemical control to attach nanocrystals to a surface [34], to electrodes [35], to other
nanocrystals, and to biomolecules [36]. Because of the relatively weak coordination
bonds of most surface-capping groups, these molecules can easily be exchanged
3.1 Semiconductor Nanoparticles
after the initial synthesis. Many of the properties of nanocrystals, notably photoluminescence, are sensitive to the surface passivation, and manipulation of the
surface-capping groups can be used to tune these properties [37, 38]. The core/
shell approach discussed in a later section proves to be a powerful method for
enhancing the fluorescence quantum yield and the stability of the particles toward
oxidation and photodegradation.
Improvement of the size distribution after the synthesis is also a widely used
strategy for achieving monodisperse samples. Size-selective precipitation is commonly applied to obtain fractions of nanocrystals with improved size distribution
from a given reaction [5, 30]. The method is based on the size-dependent solubility
of the nanocrystals, which can be used to destabilize the dispersion of the large
particles first. This is achieved by adding controlled amounts of a non-solvent to
the solution of particles, thereby leading to precipitation of the large nanocrystals,
which can be then separated by either filtration or centrifugation. Such steps can
be carried out consecutively to obtain a sequence of nanocrystal fractions with
smaller and smaller diameters. A similar procedure may in principle be applied yet
again to each fraction to further narrow the size distribution.
While the precise chemical mechanisms of nanocrystal syntheses have not been
investigated, several general comments can be made regarding the effect of the
reaction mechanism on the nanocrystal product. To accomplish the ideal nucleation event discussed above, the reaction needs to be rapid and complete. These requirements suggest simple mechanisms using reagents with low barriers to reaction. For example, at one extreme is the aqueous or micelle synthesis of CdS or
CdSe, where cadmium salts such as CdCl2 are the cadmium source and H2 S or
H2 Se is the group VI source. In these types of reactions, the metal ion is available
for reaction with a very simple group VI source. These reactions are rapid and can
even occur at room temperature. On the down side, the reactions are so facile that
it is difficult to control and isolate nucleation.
The chemistry of group III–V materials has proven to be more difficult as far as
designing a system with the attributes of both rapid and complete reaction yet
controlled nucleation. In contrast to cadmium, the group III metals are too reactive to be used as ionic sources, as they would react too strongly with the solvent
medium. Attempts to develop reactions involving group III alkyls analogous to
(CH3 )2 Cd (a reagent commonly used in the synthesis of CdSe nanocrystals) were
unsuccessful, as these compounds also appear to be too reactive, tending to proceed uncontrollably to the bulk phase regardless of capping group.
A general strategy to reduce the reactivity of these materials would be to attach
ligands to the group III and group V centers and then react these molecules, presumably in a more controlled fashion. Over the past several years, there has been
much interest in novel routes to bulk III–V materials [39–42], particularly GaAs,
because of the difficult and hazardous nature of traditional reactions involving
Ga(CH3 )3 and AsH3 . Substituted gallium and arsenic compounds were developed
in the hope of reducing reactant toxicity, eliminating gaseous reagents, and lowering reaction temperatures. As mentioned, Wells and coworkers [17] developed a
dehalosilylation reaction using GaCl3 þ As(Si(CH3 )3 )3 , which produced GaAs with
83
84
3 Syntheses and Characterizations
the elimination of (CH3 )3 SiCl. Barron et al. [43] used an analogous reaction to
produce bulk InP. These reactions are driven by the formation of the CH3 SiCl
bond in combination with the good lability of the Si(CH3 )3 ligand. They are reactions widely utilized in the synthesis of III–V semiconductor nanocrystals.
3.1.2.3
InAs and InP Nanocrystals
Synthesis of InAs and InP Nanocrystals
The synthesis of InP and InAs are analogous, and InAs will be treated in detail. All
steps of the reaction and the manipulation of the reagents are carried out under
argon or in a dry box. The setup for performing the synthesis is schematically
shown in Figure 3-13. The nanocrystals produced are crystalline and highly soluble
in a variety of organic solvents including hexanes, toluene, and pyridine. Particles
with diameters ranging from ca. 2–10 nm, with narrow size distributions, can be
prepared. The surface of the nanocrystals can be further derivatized with a variety
of ligands including amines, thiols, and phosphines.
The reaction can be written as
InCl3 þ AsðSiðCH3 Þ3 Þ3 ! InAs þ 3SiðCH3 Þ3 Cl
ð3Þ
Details of this synthesis are reported elsewhere [5, 7, 31], and here we can discuss
a typical preparation: TOP (3 g) was heated in a three-necked flask on a Schlenk
line under an Ar atmosphere to a temperature of 300 C with vigorous stirring.
Stock solution (1 mL, see below) was rapidly injected and the solution was cooled
to 260 C for further growth. The growth was monitored by taking the absorption
spectra of aliquots extracted from the reaction solution. Additional injections were
used to grow larger diameter cores. Upon reaching the desired size, the reaction
mixture was allowed to cool down to room temperature and was transferred to a
glove box. Anhydrous toluene was added to the reaction solution, and the nanocrystals were precipitated by adding anhydrous methanol. The size distribution of
the nanocrystals in a typical reaction was of the order of G15%. This was improved
using size-selective precipitation with toluene and methanol as the solvent and
non-solvent respectively.
Stock solution of InAs core: Under an Ar atmosphere, InCl3 (9 g) was dissolved in
TOP (30 mL) at 260 C with stirring. Then the solution was cooled, and taken into
the glove box. The stock solution was prepared by mixing a desired amount of
(TMS)3 As and InCl3 -TOP solution with the As:In molar ratios at 1:2 to 1:1.5.
In the stock solution, InCl3 forms a complex with TOP. While this complex has
not been isolated, several examples of InCl3 -phosphine oxide complexes have been
discussed [44], and an interaction between the Lewis acid-like InCl3 and the
strongly donating phosphine or phosphine oxide ligand would be expected. This
complex is necessary for the success of the reaction, as it helps to maintain the
solubility of the reaction intermediates and establishes the interaction that ultimately passivates the surface of the nanocrystals.
O=P
O
O
P=
O
P=
O
P=
P=O
O=P
O
P=O
P=O
P=
P=
O
P=
O
P=
room T precursors
In/GaCl3 and
P/As(SiMe3)3
O
P=
O
P = =O
P
1.
2.
Fig. 3-13 Schematic reaction setup and reaction scheme for
III–V semiconductor nanocrystals consisting of either cores
coated by organic ligands (1) or core/shell nanocrystals (2).
3.1 Semiconductor Nanoparticles
InAs + 3Me3SiCl
P
P =O
InCl3+As(SiMe3)3
liquid surfactant,
stirring and at
“high T,” 250-300 oC
85
86
3 Syntheses and Characterizations
It is valuable to examine this nanocrystal synthesis using the methodology of
nucleation, growth, and termination described earlier. While the details of the reaction’s initial stages have not been worked out, we can look to the work of Barron
[43] and Wells [45] to provide insight into the structure of these initial reaction
intermediates in the case of InP. Both groups have used the InCl3 þ P(Si(CH3 )3 )3
reaction to synthesize bulk InP at temperatures above 400 C while investigating intermediate products at lower temperatures. They report a stepwise elimination of the three equivalents of (CH3 )3 SiCl, beginning with the formation of
[Cl2 InP(SiCH3 )2 ]x at room temperature or below. The exact intermediate product
in the reaction discussed here probably has a slightly different structure due to the
action of the coordinating solvent and evidenced by the increased solubility of the
product, but the important element is the formation of the InaP bond. Wells further establishes this InaP interaction by providing a crystal structure for the related
1:1 adduct I3 InP(Si(CH3 )3 )3 . Barron sees evidence of the elimination of additional
(CH3 )3 SiCl at temperatures as low as 150 C along with the appearance of peaks in
the XPS spectrum characteristic of indium in an InP environment. However, to
eliminate the final equivalent of (CH3 )3 SiCl and obtain crystalline material, higher
temperatures are required, with reactions producing InP or InAs nanocrystals at a
minimum of about 240 C.
The termination of the particle is accomplished by the coordinating solvent medium. Both TOPO and TOP are effective in passivating the particle surface. TOP
is known as a strong donor ligand, with both high polarity and polarizability contributing to its donor strength [44, 46, 47]. Based on analogy to CdSe and standard
donor-acceptor analysis, TOP would coordinate to acceptor indium surface sites
providing a passivating shell to terminate growth, prevent agglomeration among
particles, and through its alkyl groups provide for excellent solubility in organic
solvents such as toluene and hexanes.
In order to study size-dependent phenomena, a variety of different sizes, each
with a narrow size distribution, need to be isolated. By taking advantage of the
differential solubility of various sizes of nanocrystals, narrow size distributions can
be obtained using size-selective precipitation techniques. The solubility of a specific nanocrystal is determined by several factors, most notably the size of the particle, its shape, and the nature of the particle surface. Because of the alkyl groups
of the TOP cap, the particles are soluble in relatively nonpolar solvents. If a more
polar solvent is added to a solution of nanocrystals, the ability of the combined
solvent system to solvate the nanocrystals will be reduced. The mechanism of
agglomeration is dominated by attractive Van der Waals forces between particles
[48–51], and at some composition of solvents the solvation ability of the system
will no longer overcome these attractive forces and the particles will begin to agglomerate and precipitate. The key to size-selective precipitation is the fact that
different sizes, shapes, and surface coverages will have different solubilities and
thus can be separated. Assuming that the dominant difference in the nanocrystals
is that of particle size, the ensemble of nanocrystals will be separated by size, while
the degree of resolution among sizes will be determined by the smaller differences
in the sample such as shape or surface structure. The largest particles will precip-
3.1 Semiconductor Nanoparticles
itate first because of their stronger attraction. One primary solvent system employed is toluene/methanol. The nanocrystals are soluble in toluene, but are insoluble in highly polar methanol. Starting with the initial reaction mixture diluted in
toluene, the incremental addition of methanol results in the size-selective precipitation of the nanocrystals and the isolation of nanocrystals with individual size
distributions.
Structural and Basic Optical Characterization of InAs and InP Nanocrystals
The evaluation of nanocrystal structure encompasses several important areas including the identity and crystallinity of the core, the morphology of the particles,
and the composition of the nanocrystal surface. The crystallinity of the core can be
studied using powder XRD, electron diffraction, and the morphology of the particles, including their shape and size distributions, determined by transmission
electron microscopy (TEM). X-ray photoelectron spectroscopy (XPS) can be used to
examine the nanocrystal composition, with particular attention to the surface capping groups and will be discussed in a later section concerning the core/shell
nanocrystals. InAs and InP nanocrystals will be treated together, with representative data shown from one or both systems.
Powder X-ray diffraction spectra of a series of InAs nanocrystal sizes are shown
in Figure 3-14. The peak positions index well with the bulk InAs lattice reflections.
Information on the domain size of the sample can be obtained from the width of
the lattice reflections using the Debye-Scherrer formula for spherical particles [52]:
D¼
1:2l
½ðD2yÞðcos yÞ
In this equation, l is 1.5418 Å for CuKa radiation, D2y is the measured linewidth
in radians, y is 1/2 of the measured diffraction angle, and D is the estimated particle size in Angstroms. An average domain size for a given sample was obtained
by averaging D for the (111), (220), and (311) peaks and is shown in the figure.
TEM is crucial for investigating the structure of the nanocrystals as it allows a
direct evaluation of both the size and shape of the particles. A high-resolution TEM
image of InAs nanocrystals, 2 and 6 nm in diameter, is shown in Figure 3-15.
Cross fringes are observed, indicating a high degree of crystallinity, with the lattice
spacings indexing to those of InAs. The nanocrystals are almost spherical.
TEM is also essential for calculating the size and size distribution of a nanocrystal sample. For the wide variety of size-dependent properties exhibited by the
nanocrystals, this is clearly a crucial measurement. Size distributions can be obtained by the direct measurement of images of large numbers of nanocrystals,
typically a few hundred. The results of these measurements for InP show size distributions with standard deviations of about 20%, while for InAs the distributions
are substantially narrower than 10%.
The quality of the size distribution is most clearly borne out in Figure 3-16,
showing ordered self-assembled solids of InAs nanocrystals, with diameter of 5.8
87
88
3 Syntheses and Characterizations
Powder X-ray diffraction spectra of InAs nanocrystals
of different sizes. The domain size was calculated from the
width of the reflections as detailed in the text. Reproduced
from Appl. Phys. Lett. 69, 1432 (1996).
Fig. 3-14
nm, dissolved in toluene, and deposited onto the TEM grid. Figures 3-16A–C exhibit the real space images for three different superlattice (SL) faces detected on
the grid, while Figures 3-16D–F represent the optical diffraction patterns measured for the corresponding TEM negatives. Ordered domains extending over regions of lengths measured in microns were observed. Both the real space images
and the diffraction patterns, which in effect correspond to a two-dimensional
Fourier transform of the image, allow the assignment of the three different faces to
different views along three zone axes of the close-packed face centered cubic (FCC)
structure. Figure 3-16A and its corresponding diffraction pattern Figure 3-16D
represent the superlattice when viewed along the (111)SL zone axis, Figure 3-16B
(diffraction Figure 3-16E) represent the (110)SL zone axis, and Figure 3-16C (diffraction Figure 3-16F) represent the (100)SL zone axis. Figure 3-16G–I show images of ordered monolayers with three different packings corresponding to three
different superlattice faces resolved in Figures 3-16A–C respectively.
3.1 Semiconductor Nanoparticles
Fig. 3-15 HRTEM images of two InAs nanocrystals, 2.2 and
5 nm in diameter. Reproduced from Appl. Phys. Lett. 69, 1432
(1996).
Quantum confinement effects of nanocrystals are evidenced most clearly in the
optical properties of the system as the electronic energy levels of the clusters become a function of size. A detailed account of this aspect will be provided in
Chapter 5. Basic optical characterization of semiconductor nanocrystals provides
important information on particle size, from the position of the band gap energy,
and on the size distribution, from the sharpness of peaks in absorption and luminescence. Figure 3-17 shows room temperature absorption spectra for a series of
InAs nanocrystal sizes along with the photoluminescence spectra. The quantum
confinement effects are clearly evident from the size-dependent nature of the
spectra. The band gap in all samples is shifted substantially from the bulk InAs
gap of 0.42 eV. In all samples, the onset of absorption is characterized by a distinct
feature at the absorption edge corresponding to the first excitonic transition. Additional features are resolved at higher energy, corresponding to higher excited
states. All states shift with nanocrystal size. The width of the features is primarily
due to inhomogeneous broadening resulting from size variations of the particles
within a specific sample. The size distribution as obtained from TEM is shown in
the right-hand frame of Figure 3-17. The absorption of InP nanocrystals shows
similar effects, the absorption onset shifting to the blue with decreasing cluster
diameter as discussed in [4].
III–V Core/Shell Nanocrystals – Synthesis and Characterization
The emission color from semiconductor nanocrystal quantum dots is tunable by
the size as a result of the quantum-confinement effect [53, 54]. Harnessing this
emission for real-world applications such as biological fluorescence marking [55,
56] and optoelectronic devices [11, 57–60] is an important challenge, which imposes stringent requirements for a high fluorescence quantum yield (QY) and stability against photo-degradation. These characteristics are difficult to achieve in
3.1.2.4
89
90
3 Syntheses and Characterizations
A
B
D
E
G
F
H
TEM images of arrays and superlattices prepared from InAs nanocrystals,
5.8 nm in diameter. Frames A–C show ordered
superlattices prepared from InAs nanocrystals,
Fig. 3-16
C
I
while frames D–F show the optical diffraction
pattern from the micrographs. Frames G–I
show TEM images of monolayers. See text for
details.
semiconductor nanocrystals coated with organic ligands because of imperfect surface passivation. In addition, the organic ligands are labile in exchange reactions
because of their weak bonding to the nanocrystal surface atoms [37]. As discussed
previously, a proven strategy for increasing both the fluorescence QY and the stability is to grow a shell of a higher band gap semiconductor on the core nanocrystal
[23–27, 61]. In such composite core/shell structures, the shell type and shell
thickness provide further control for tailoring the optical, electronic, electrical and
chemical properties of semiconductor nanocrystals. Several reports of III–V core/
shell nanocrystals have already been mentioned, and here we shall focus on our
own detailed study of core/shell nanocrystals with InAs cores [6–7]. The band gap
energy of these core/shell nanocrystals is tunable in the near IR (NIR) spectral
range, covering the wavelengths that are important for telecommunication applications. Recently, they were successfully incorporated with semiconducting polymers to form an LED structure that provides efficient electroluminescence at
3.1 Semiconductor Nanoparticles
Absorbance and Emission (a.u.)
28 Å
0.8
43 Å
50 Å
60 Å
1.4
1.8
2.2
17
Energy [eV]
The right-hand panel shows the
distribution of diameters for 4 InAs nanocrystal
samples obtained from HRTEM measurements. The left-hand panel displays the room
temperature absorption (solid lines), and
Fig. 3-17
45
73
Diameter [Å]
emission spectra (dashed lines) for the
corresponding samples. The mean diameter
of each sample is noted. Reproduced from
Superlatt. Microstruct. 22, 559 (1996).
1.3 mm [11]. These core/shell nanocrystals may be further developed to serve as
efficient fluorescent labels for biological applications in the NIR range.
The core/shell approach is conceptually closely related to the approach used in
two-dimensional quantum wells [62]. In a quantum well, a thin layer, of nanometric dimensions, of a low band gap semiconductor is sandwiched between thick
layers of a high band gap semiconductor, forming a square potential well for the
electron and hole wavefunctions. The determining factor for growth of quantum
wells is the lattice mismatch between the two semiconductors. The electronic
properties of the quantum well are determined primarily by the energetic offsets
between the conduction and valence band edges of the two semiconductors.
Therefore, the influence of the lattice mismatch and the band offsets on the growth
and electronic properties of core/shell nanocrystals with InAs cores was investigated. Figure 3-18 shows schematically the values of both parameters for InAs
and a variety of shell materials [63]. The lattice mismatch ranges from nearly zero
for CdSe shells to as high as 10.7% for ZnS shells. The band offsets cover a broad
range of values as well. Note that the shell materials include both II–VI and III–V
91
92
3 Syntheses and Characterizations
Shell
In As
1.74
CB
1.26
0.94
0.51
0.46
GaAs
In P InAs
0.06
CdS ZnS ZnS
e
e
0.42
6.67%
VB
0.92
3.13%
0.99
0.00%
Lat
1.52
tice +0.001%
Mis
6.44%
ma
10.7%
tch
Summary of the band offsets (in eV)
and lattice mismatch (in %) between the core
InAs and the III–V semiconductor shells
(left side) and II–VI semiconductor shells
Fig. 3-18
(right side) grown in this work. CB ¼
conduction band; VB ¼ valence band.
Reproduced from J. Am. Chem. Soc. 122,
9692–9702 (2000).
semiconductors, a selection that provides substantial tunability for the properties
of the composite core/shell nanocrystals.
Synthesis of Core/Shell Nanocrystals with InAs Cores
The preparation of the InAs core/shell nanocrystals is carried out in a two-step
process. In the first step the InAs cores are prepared using the injection method
with TOP as solvent, which has enabled us to obtain hundreds of mg of nanocrystals per synthesis, as detailed earlier. We used size-selective precipitation to
improve the size distribution of cores to s @ 10%. In the second step, the shells of
the various materials were grown on these cores. This two-step approach follows
the methods developed for synthesis of the II–VI core/shell nanocrystals.
The details of the synthesis of the core/shell nanocrystals can be found elsewhere [6, 7]. Briefly, TOP-capped InAs cores (5–20 mg) were dissolved in TOP
(3–6 g, or a mixture with TOPO) in a three-necked flask. Under Ar flow on a
Schlenk line, the nanocrystal solution was heated to 260 C, and the shell precursor solution was introduced into the hot solution by drop-wise addition. The
growth of core/shells was monitored by UV-VIS spectroscopy of aliquots taken
from the reaction flask. After the desired shell thickness had grown, the reaction
3.1 Semiconductor Nanoparticles
mixture was cooled to room temperature. Core/shell nanocrystals passivated by
TOP were obtained by precipitation using a mixture of methanol and toluene.
Attempts to carry out the growth of III–V semiconductor shells, InP and GaAs,
at a temperature of 160 C in TOP were not successful. Low temperature reduces
the rate of nucleation of nanocrystals of the shell material, which is an undesired
process competing with shell growth. In addition, the alloying process of core and
shell materials is less likely at low temperatures. However, growth could not be
achieved at this low temperature. These shells could be grown only at higher temperatures: T > 240 C. This minimal temperature is required to overcome the
reaction barrier for the precursors, similarly to the conditions required for growth
of III–V semiconductor nanocrystal cores. Above this limit, controlled growth of
InP shells of varying thickness was achieved.
Thin GaAs shells could also be grown, but, in contrast to InP shells, growth was
limited to a thickness of less then two monolayers. This difference may be related
to strong bonding of Ga to TOP, consistent with the observation that it is difficult
to synthesize GaAs nanocrystals in this way [64]. An additional difference between
the GaAs and the InP shells is in the solubility of the core/shell nanocrystals of
each kind. InAs/InP core/shells are readily soluble in organic solvents after precipitation of the nanocrystals from the growth solution. Special conditions were
required to obtain good solubility of the InAs/GaAs core/shells [1].
In contrast to the III–V semiconductor shells, the growth of the II–VI semiconductor shells, CdSe, ZnSe and ZnS, was observed even at much lower temperatures (above 150 C). For CdSe shells, it was found that to limit the nucleation
and achieve controlled shell growth TOP could be used instead of TOPO as the
growth medium. This is consistent with stronger bonding of Cd to TOP versus
TOPO.
For InAs/ZnSe and InAs/ZnS core/shells grown in TOP, solubility after separation from the growth solution is poor as the TOP ligands can easily be removed by
methanol. Additionally, upon growth of thick ZnS shells, substantial nucleation of
ZnS nanocrystals was observed by XRD measurements. To overcome these difficulties, a TOP-TOPO mixture can be used as the growth solution instead of TOP.
In this case, nucleation could be minimized, and soluble nanocrystals were obtained. The growth rate of ZnSe and ZnS shells was substantially lower for the
TOP-TOPO mixture compared with TOP. These features are consistent with
stronger bonding of Zn to TOPO versus TOP, as was reported in the synthesis of
ZnSe nanocrystals [65]. Cd is a softer Lewis acid than Zn and binds strongly to
TOP, a softer Lewis base than TOPO, while Zn, the harder Lewis acid, binds more
strongly to TOPO the harder Lewis base. This follows the general hard-soft concept
for interaction strength between Lewis acids and bases [66].
Optical Characterization of the Core/Shell Nanocrystals
The most direct and immediate probes for shell growth are the absorption and
fluorescence spectra. Figure 3-19 shows the sequence of absorption spectra measured for aliquots taken from the reaction solution during the growth of InP shells
on InAs cores with an initial radius of 1.3 nm. The first absorption peak shifts to
93
3 Syntheses and Characterizations
1.2
d
933 nm
886 nm
991nm
e
Absorbance (a.u.)
94
1.4
1.6
1.8
2.0
c
b
a
2.2
Photon Energy (eV)
Evolution of the absorption spectra during the
growth of InP shells on InAs cores with an initial radius of
1.3 nm (a). The InP shell thickness in number of monolayers is:
(b) 0.5, (c) 1.1, (d) 1.7, (e) 2.5. Reproduced from J. Am. Chem.
Soc. 122, 9692–9702 (2000).
Fig. 3-19
the red upon shell growth. The red shift occurs because the conduction band offset
between InAs and InP is smaller than the confinement energy of the electron, and,
as the shell grows, the electron wavefunction extends to a larger box and its confinement energy is lowered. As the electron effective mass, me , in InAs is extremely small (me ¼ 0:024m e , m e is the mass of the free electron [67]), it is highly
delocalized, and a large potential step is required to confine it to the core. The
spectral features remain sharp during the reaction, indicating that the growth is
controlled and the size distribution is maintained. However, the fluorescence of the
InAs/InP core/shells is quenched substantially as compared to the original cores.
For CdSe, which has a conduction band offset similar to that of InP, shell growth
also leads to a red shift of the absorption onset. Figure 3-20A,B shows the evolution
of the absorption and the emission spectra during growth of InAs/CdSe core/
shells with two different initial core radii, 1.2 nm (Figure 3-20A), and 2.5 nm
(Figure 3-20B). As in the case of InP shells, this red shift results from the lower
confinement energy of the electron, whose wavefunction extends to the shell region. In this case, the sharpness of the spectral features is partially washed out
during the growth. In contrast to InAs/InP core/shells, the band gap fluorescence
QY for InAs/CdSe core/shells is substantially enhanced, up to a maximum value
of 21%, nearly 20 times larger than the QY of the cores. These values of the QY
compete favorably with QY values for organic NIR laser dyes [12]. This pronounced
difference in emission of InP versus CdSe shells was ascribed to the different in
the quality of the outer surface of the nanocrystals. In both cases, the similar and
relatively small conduction band offset leads to a substantial probability of the
presence of the electron wavefunction at the nanocrystal surface. As a result, for
3.1 Semiconductor Nanoparticles
B)
D)
C)
F)
E)
Absorbance (a.u.)
Photoluminescence Intensity (a.u.)
A)
95
x2
x4
x2
x4
x2
0.8 1.2 1.6 2.0 2.4 2.8 0.8 1.2 1.6 2.0 2.41.2 1.6 2.0 2.4 2.8 0.8 1.2 1.6 2.0 2.4 1.2 1.6 2.0 2.4 2.8
Energy (eV)
Energy (eV)
Energy (eV)
Evolution of absorption (dashed
lines), and photoluminescence (solid lines)
for growth of core/shells. The PL spectra are
given on a relative scale for comparison of
the enhancement of QY with shell growth.
Reproduced from J. Am. Chem. Soc. 122,
9692–9702 (2000). A. InAs/CdSe with initial
core radius of 1.2 nm. The shell thickness
(in number of monolayers) and QY for the
traces from bottom to top are respectively :
0, 1.2%; 0.6, 13%; 1.2, 21%; 1.8, 18%. B.
InAs/CdSe with initial core radius of 2.5 nm.
The shell thickness (in number of monolayers)
and QY for the traces from bottom to top are
respectively : 0, 0.9%; 0.7, 11%; 1.2, 17%; 1.6,
14%. C. InAs/ZnSe with initial core radius of
1.2 nm. The shell thickness (in number of
Fig. 3-20
Energy (eV)
Energy (eV)
1.2 1.5 1.8 2.1 2.4
Energy (eV)
monolayers) and QY for the traces from
bottom to top are respectively : 0, 1.2%; 0.6,
9%; 1.5, 18%; 2.5, 14%. D. InAs/ZnSe with
initial core radius of 2.8 nm. The shell thickness (in number of monolayers) and QY for
the traces from bottom to top are respectively :
0, 0.9%; 0.7, 13%; 1.3, 20%; 2.2, 15%. E.
InAs/ZnS with initial core radius of 1.2 nm.
The shell thickness (in number of monolayers)
and QY for the traces from bottom to top are
respectively : 0, 1.2%; 0.7, 4%; 1.3, 8%; 1.8,
7%. F. InAs/ZnS with initial core radius of
1.7 nm. The shell thickness (in number of
monolayers) and QY for the traces from
bottom to top are respectively : 0, 1.1%; 0.6,
5%; 1.3, 7.1%; 2.2, 6.3%.
both cases the emission is sensitive to the outer surface and susceptible to trapping
in unpassivated surface sites. Indeed, this problem is known for InP nanocrystals
prepared in TOP/TOPO, which have a very low fluorescence QY that could be enhanced by surface treatments such as oxidation and etching [22]. CdSe TOP/
TOPO-coated nanocrystals, on the other hand, display band gap emission with a
room temperature QY of a few percent.
ZnSe and ZnS are shell materials with much larger band offsets relative to that
of InAs. The evolution of the absorption and emission spectra for ZnSe and ZnS
shell growth on cores with various radii is presented in Figures 3-20C,D and 3-20E,F
respectively. For both core/shells, the absorption onset is nearly the same as in the
original core. This can be explained by the large conduction and valence band offsets of ZnSe and ZnS relative to InAs, which create a substantial potential barrier
for the electron and hole wavefunctions at the core/shell interface. Unlike the CdSe
and InP shells, here the electron and hole wavefunctions are both confined primarily to the core region, and thus the band gap does not shift from the value of
96
3 Syntheses and Characterizations
the core. As shown in Figures 3-20C–F, the fluorescence QY is enhanced for both
shells. For ZnSe the maximum QY values are 20%, while in ZnS the maximum
QY is 8%, achieved for a shell thickness of 1.2–1.8 monolayers. With further shell
growth, the QY decreases. This reduction may be ascribed to trapping of the charge
carriers at the core/shell interface. The lattice mismatch between InAs and both
ZnSe and ZnS is large. In thin shells the strain can still be sustained, and epitaxial
growth of the shell on the core can occur. With further shell growth, defects may
form at the core/shell interface, and these could trap the carriers, leading to the
gradual reduction of QY in the thicker shells. A similar observation was reported
for the II–VI semiconductor core/shells CdSe/ZnS and CdSe/CdS [25, 26].
The selection of core/shells extends the control afforded over the electronic
properties of the composite nanocrystals. As an example of the flexible control, two
types of core/shells are demonstrated, InAs/ZnSe and InAs/CdSe, which emit
strongly at 1.3 mm, a wavelength with significant technological importance in fiberoptic communications. Figure 3-20B presents the absorption and emission for a
CdSe shell overgrown on a core with a radius of 2.5 nm. The core band gap emission is at 1220 nm, and with the growth of the shell the emission shifts to the red.
This is accompanied by substantial enhancement of the QY, up to a value of 17%
achieved at 1306 nm. For ZnSe shells the band gap hardly shifts. Figure 3-15D
shows that using a bigger core, with radius of 2.8 nm, we could achieve a high QY
of 20% at 1298 nm by growing the ZnSe shell. Such core/shell nanocrystals were
recently used as the optically active chromophores in efficient light-emitting diodes
covering the near IR spectral range [11]. The LED was based on a nanocrystalsemiconductor polymer composite.
Chemical and Structural Characterization
X-ray photoelectron spectroscopy (XPS): XPS can be used effectively to examine
the chemical composition of the core/shells [68]. Figure 3-21 shows the XPS survey spectra for InAs cores of radius 1.7 nm and for InAs/CdSe core/shells prepared with the same core and with shell thickness of 3 monolayers. Indium and
arsenic peaks are clearly resolved for the core (lower spectrum). Additional peaks
belonging to Cd and Se, the shell materials, can be identified in the core/shells
(top spectrum). The ratio of the core atom XPS peak heights for cores to core atom
XPS peak heights for core/shells is energy dependent. The relative intensity of the
peaks at high binding energies (e.g., InMNN ), which have a small kinetic energy
and thus a small escape depth, is quenched more upon shell growth, compared
with the peaks at small binding energy and large escape depth (e.g., AsLMN ).
High-resolution XPS provides quantitative evidence for shell growth based on
the finite escape depth, l, of photoelectrons from the core atoms [26]. The typical
escape depths are of the order of the shell thickness, and the photoelectron signal
from core atoms should decrease accordingly in the core/shell structure [68].
The application of XPS to the study of shell growth is demonstrated in Figure
3-22, showing the XPS data for InAs/ZnSe core/shells with an initial core radius of
1.7 nm. We performed high-resolution XPS measurements on a sequence of core/
100
200
600
* Cd MNN
In 3d * Cd 3d3
O 1s
* Cd 3p3
* Cd 3p1
In 3p3
In 3p1
* Cd 3d5
C 1s
As LMN
Au 4d5
Au 4d3
300
In MNN
As 3s
As 3p
0
As LMN
C 1s
Au 4d5
Au 4d3
In 3d5
O 1s
In 3d3
Au 4p3
In 3p3
In 3p1
Au 4f7
Au 4f5
* Se LMN
* Se 3p
* Se LMM
As 3s
* Cd 4d
In 4d
As 3d
As 3p
In 4d
As 3d
XPS Intensity (a.u.)
3.1 Semiconductor Nanoparticles
900
1200
Binding energy(eV)
XPS survey spectrum for InAs cores
with radius of 1.7 nm (lower trace), and for
InAs/CdSe core/shells with shell thickness of 3
monolayers (top trace). The assignment of the
peaks is indicated. The new peaks in the XPS
Fig. 3-21
spectrum for the core/shells associated with
Cd and Se are emphasized in bold italic type.
Reproduced from J. Am. Chem. Soc. 122, 9692–
9702 (2000).
shell samples of varying thickness. The figure shows the experimental results
(squares) for the log of the ratio of the In3d5 to the InMNN Auger peak, normalized
by the ratio in the core. The ratio increases upon ZnSe shell growth. The number
of atoms for the two peaks is identical, and the increase in the ratios is due to the
difference in the escape depths, which is lower for the Auger peak. As a result,
the shell growth leads to larger reduction in the relative intensity of the Auger
peak. To check these effects, the expected ratio for a spherical geometry was simulated assuming a core/shell configuration (solid line), or alloy formation (dashed
line). Clearly, only the calculated ratio for the core/shell structure is in agreement
with experimental data.
Transmission electron microscopy : Figure 3-23 presents HRTEM images of core
and core/shell nanocrystals. Frame B shows a micrograph of an InAs core with a
97
3 Syntheses and Characterizations
0.16
XPS : log( In3d5 : InMNN)
98
(InAs)/(ZnSe)
0.12
0.08
0.04
0.00
0
5
10
15
Shell Thickness (Å)
Fig. 3-22
Summary of high-resolution XPS
data for InAs/ZnSe. The log of the ratio of the
intensity of the In3d5 to that of the InMNN is
plotted against the shell thickness. The ratio is
normalized to the ratio in the cores. Squares:
experimental data. Solid line: calculated ratio
for core/shell structure. Dotted line: calculated
ratio for alloy formation. Reproduced from
J. Am. Chem. Soc. 122, 9692–9702 (2000).
HRTEM images of InAs/InP core/
shell (frame A, core radius 1.7 nm, shell
thickness 2.5 nm), InAs core (frame B, core
radius 1.7 nm), and InAs/CdSe core/shell
(frame C, core radius 1.7 nm, shell thickness 1.5 nm). The scale bar is 2 nm. The
nanocrystals are viewed along the [011] zone
axis. Frame D shows the Fourier transform of
image C, and the pattern corresponds to the
diffraction pattern from the 011 zone of the
cubic crystal structure. Reproduced from
J. Am. Chem. Soc. 122, 9692–9702 (2000).
Fig. 3-23
radius of 1.7 nm, and frames A and C respectively show InAs/InP and InAs/CdSe
core/shell nanocrystals with thick shells grown on similar cores. The crystalline
interior is viewed along the [011] zone axis of the cubic lattice. The cubic lattice
is resolved for the InAs/CdSe core/shell as well, as clearly revealed also by the
Fourier transform of the image presented in Figure 3-23D. CdSe nanocrystals
grown in such conditions form the wurtzite structure, but here they adopt the
cubic structure on the InAs core. For all three cases, the fringes are visible across
the entire nanocrystals in accordance with epitaxial shell growth in these particles.
Intensity (a.u.)
Intensity (a.u.)
3.1 Semiconductor Nanoparticles
Peak 111 Zoom
20
24
28
32
2
c
b
10
20
30
40
50
60
70
80
531
440
90
440
531
511
422
511
422
400
400
331
420
331
420
311
311
222
222
220
220
200
200
Zinc Blende:
InAs
InP
111
111
a
100
2q
XRD patterns for InAs cores with
radius of 1.7 nm (trace a, solid line), and InAs/
InP core/shells with shell thickness of 2.4
monolayers (trace b, dashed line), and 6.2
monolayers (trace c, dot-dashed line). The
Fig. 3-24
inset shows a zoom of the 111 peak. The
diffraction peak positions of bulk cubic InAs
(dark) and InP (light, italic) are indicated.
Reproduced from J. Am. Chem. Soc. 122, 9692–
9702 (2000).
X-ray diffraction: Powder XRD patterns for the InAs core, radius 1.7 nm, and for
InAs/InP core/shells with increasing thickness are presented in Figure 3-24. The
InAs XRD pattern consists of the characteristic peaks of cubic InAs, which are
broadened because of the finite crystalline domain. With InP shell growth, the
diffraction peaks shift to larger angles consistent with the smaller lattice constant
for InP compared with InAs. The shift is most clearly resolved for the high angle
peaks. In addition, the diffraction peaks narrow. This is demonstrated for the (111)
peak shown in the inset of Figure 3-24. This narrowing indicates that the crystalline domain is larger for the core/shells, providing direct evidence for epitaxial
growth mode of the shell. The relatively simple diffraction pattern for cubic InAs
allows a clear observation of the narrowing as the shell is grown. In contrast, in the
case of cores/shells with CdSe cores, the narrowing is masked by the complex diffraction pattern of the wurtzite structure [15].
The XRD patterns for a series of core/shells with different shell materials and
varying thickness are displayed in Figure 3-25 (filled circles). The general pattern
99
3 Syntheses and Characterizations
InAs/ZnS
InAs/GaAs
InAs/ZnSe
Intensity (a.u.)
100
InAs/InP
InAs/CdSe
InAs
111
10
20
200
30
220
40
311
50
60
2q
X-ray diffraction patterns for InAs
cores and various core shells. The experimental
curves (filled circles) are compared with the
simulated curves (thin solid lines). See text
and Table 1 for the details of the simulated
structures. The markers on the bottom of each
frame indicate the diffraction peak positions
Fig. 3-25
and the relative intensities for the InAs core
material (lower frame) and the various shell
materials (other traces). The vertical dashed
lines indicate the positions of the InAs core
nanocrystal diffraction peaks. Reproduced from
J. Am. Chem. Soc. 122, 9692–9702 (2000).
3.1 Semiconductor Nanoparticles
of the cubic lattice is maintained for all materials. The diffraction peaks narrow
with shell growth in the case of CdSe, ZnSe, and GaAs shells for the reasons discussed above. But this is not the case for ZnS, probably because of the very large
lattice mismatch of ZnS and InAs (10.7%). In this case, the large strain may lead
to cracking at the InAs/ZnS interface already at early growth stages. Moreover,
ZnS nanocrystals in these growth conditions preferentially form the wurtzite
structure, which may create additional defects upon shell growth. Both effects may
explain the relatively small enhancement of the PL QY for the ZnS shells.
To further prove the interpretation of the XRD data and to obtain more quantitative information about the core/shell structures, the powder diffraction patterns
were simulated [7, 69]. Nanocrystals are built by stacking planes along the (111)
axis of the cubic lattice. The sum of the specified core radius, rc , and shell thickness, rs , was used to carve out the nanocrystal assuming a spherical shape.
The experimental data (filled circles) and the simulation results (thin lines) for
the core and a series of core/shells are displayed in Figure 3-25. The simulation of
the cores fits the experimental data very well. The fit was obtained using the simulated XRD pattern for an equally weighted combination of two kinds of core
nanocrystal structures with the same radii, differing in the (111) plane stacking
sequence. Both nanocrystal structures have three stacking faults. The experimental
diffraction intensity between the (220) and the (311) peaks does not go to zero,
while in the simulated pattern for nanocrystals without stacking faults the value is
close to zero. Thus, three stacking faults along the (111) direction are required to
quantitatively reproduce the experimental pattern. Introduction of surface disorder
has little effect on the simulated patterns.
For InAs/CdSe core/shells, the lattice mismatch is zero. The experimental peak
positions do not shift with shell growth. This is well reproduced in the simulation.
An additional stacking fault was added in the shell region for the thicker shell of
three monolayers, to better reproduce the experimental pattern.
For the other core/shells there is a lattice mismatch, which ranges from 3.13%
for InP to 10.7% for ZnS. A gradual shift of the diffraction peaks with shell growth
towards larger angles is observed for these core/shells (Figures 12 and 13). This
means that the lattice spacing is modified in the shell region. In order to achieve
epitaxial growth mode, atoms in both core and shell regions at the core/shell interface must have identical lattice spacings. To simulate the smooth switching of
the lattice spacing from the core to the shell, a Fermi-like switching function was
used. This switching provides a physical model for understanding the epitaxial
growth mode, while reproducing the observed change in the peak positions upon
shell growth.
The simulated patterns reproduce the two main effects observed upon shell
growth in InAs/InP, InAs/ZnSe, and InAs/GaAs core/shells, namely the shifting
of all the diffraction peaks to larger angles and the narrowing of the peaks (Figure
3-25). This provides further evidence for the epitaxial shell growth mode. For InAs/
ZnS, no narrowing was observed in the experimental diffraction pattern, indicating
that in InAs/ZnS core/shells the interface region is not fully epitaxial, as explained
above. Similarly to the case of InAs/CdSe core/shells, for the thick InP shells (6.2
101
102
3 Syntheses and Characterizations
monolayers), stacking faults had to be added in the shell region to better reproduce
the experimental pattern. On average, a stacking fault exists for every four to five
layers in InAs core/shell nanocrystals. In these stacking faults the bonds of atoms
remain fully saturated, and charge carrier traps do not necessarily form. Therefore,
it is likely that these planar stacking faults do not reduce the fluorescence QY
substantially.
Model Calculations for the Band Gap
The core/shell band offsets provide control for modifying the electronic and optical
properties of these composite nanocrystals. To examine the effect of the band offsets of various shells on the band gap of the composite nanocrystals, calculations
using a particle-in-a-spherical-box model were performed [16, 70]. Briefly, in this
model the electron and hole wavefunctions are treated separately, and the coulomb
interaction is then added within first order perturbation theory [71]. Three radial
potential regions should be considered in the core/shell nanocrystals: core, shell
and the surrounding organic layer. Continuity is required for the radial part of the
wavefunctions for both electron and hole at the interfaces. In addition, the probability current, ð1=m i Þðd=drÞR i ðk i rÞ, where m i is the effective mass in region i, R i
is the radial part of the lowest energy 1Se=h electron or hole wavefunction, and k i is
the wave vector in region i, has to be continuous. The effective masses and dielectric constants of the bulk semiconductors were used in the calculations [67]. The
band offsets were taken from [63]. First, for InAs cores of radius 1.7 nm, a barrier
height of 4.5 eV was used for the surrounding organic layer for both carriers. The
confinement energy for the electron is sensitive to the barrier height because of the
small electron effective mass, while the heavier hole is much less sensitive. The
experimental value is in reasonable agreement with the calculated band gap, although there is a deviation, which is understandable in this simplistic model [72].
The calculated versus measured relative change of the band gap upon growth of
the different shells is plotted for various core/shells in Figure 3-26. The error in the
band gap shifts should be small compared with that in the absolute band gap
energies. The agreement for the band gap shifts is good. For the ZnS shells (dotdashed line), hardly any shift is observed either in the experiment or in the calculations. This is consistent with the large band offsets between InAs and ZnS. For
both InP (dashed line) and CdSe (solid line) shells, the experimental and calculated
band gaps shift to the red upon shell growth. This is mainly because of the reduction in the electron confinement energy in these core/shells. The lowest 1Se level
of the electron is above the core/shell barrier for both InP and CdSe shells, leading
to a high probability of the presence of the electron in the shells.
Stability of Core/Shell Nanocrystals
Core/shell passivation with a shell of a semiconductor material that has large band
offsets compared with the core also provides increased protection and stability
compared with the organically passivated core nanocrystals. The photostability of
the core/shells was compared to IR140, a typical near IR organic laser dye [73].
Solutions of nanocrystals and of the dye saturated with oxygen were irradiated at
3.1 Semiconductor Nanoparticles
Ecore/shell - Ecore (m eV)
0
-50
-100
-150
-200
-250
0
5
10
15
20
25
Shell thickness(Å)
Experimental and calculated shifts of
the band gap energy in various core/shells
versus the shell thickness. InAs/ZnS core/
shells: experimental data – squares, calculated
shift of gap – dot-dashed line. InAs/InP core/
Fig. 3-26
shells: experimental data – circles, calculated
shift of gap – dashed line. InAs/CdSe core/
shells: experimental data – triangles, calculated
shift of gap – solid line. Reproduced from
J. Am. Chem. Soc. 122, 9692–9702 (2000).
810 nm with a Ti-sapphire laser for 30 min. An intensity of 250 mW was used, and
the optical density of the irradiated solutions was 0.2 at 810 nm. Under these
conditions, each nanocrystal and dye molecule absorbs a total of approximately
0:5 10 6 photons. For the dye, following the irradiation, the main absorption peak
at 830 nm vanished, indicating that the dye had completely degraded. For InAs
cores (radius 1.7 nm), the absorption maximum shifted to the blue by 10 nm and
the optical density slightly decreased. The PL intensity decreased by a factor of 2.1.
For InAs/ZnSe core/shells (core radius 1.7 nm, shell thickness @2 monolayers),
the absorption did not change upon the irradiation, and the PL intensity decreased
by a factor of 1.7 to a value still eight times stronger than that for the fresh core.
Finally, for InAs/CdSe core/shells (core radius 1.7 nm, shell thickness @1.8 monolayers), the absorption shifted to the blue by 5 nm, and the PL intensity decreased
by only 10%. This experiment demonstrates the improved photostability of the
core/shells compared with a typical near-IR laser dye.
These core/shells also display greatly increased stability toward oxidation compared with the cores. The stability of the bare (TOP-coated) InAs cores and that of
the two core/shell samples similar to the ones examined above were studied. Absorption and emission spectra of the fresh samples and of samples that were kept
for ten months in a solution saturated with oxygen under daylight conditions were
compared, as summarized in Figure 3-27. The InAs cores under these conditions
103
3 Syntheses and Characterizations
A: InAs core
x 20
~ 0.03%
~1.2%
B: InAs/CdSe
Absorbance (a.u.)
Photoluminescence Intensity (a.u.)
104
~ 13%
~ 16%
C: InAs/ZnSe
~ 15%
~ 19%
0.9
1.2
1.5
1.8
2.1
2.4
Photon Energy (eV)
Comparison of stability of cores
(frame A), InAs/CdSe core/shells (frame B),
and InAs/ZnSe core/shells (frame C). Absorption spectra (solid lines) and PL spectra
(dashed lines) are shown on the same energy
scale for fresh nanocrystals (bottom traces in
Fig. 3-27
each frame) and for nanocrystals kept in an
oxygen-saturated solution for 10 months (top
traces in each frame). The photoluminescence
QY (%) is also indicated in each case.
Reproduced from J. Am. Chem. Soc. 122, 9692–
9702 (2000).
3.1 Semiconductor Nanoparticles
exhibited a considerable blue shift of the absorption accompanied by washing out
of the spectral features. The emission was quenched by a factor of 40 compared
with the fresh cores (Figure 3-27A). These phenomena indicate substantial oxidation of the bare cores. The core/shells under similar conditions show a very different behavior. For InAs/CdSe core/shells the absorption shifted slightly to the blue
and the QY decreased from 16% to 13% (Figure 3-27B). Finally, for the InAs/ZnSe
core/shells, the absorption hardly shifted and the emission QY was reduced from
19% to 15% (Figure 3-27C).
Acknowledgement
I wish to thank my coworkers and colleagues who have contributed immensely to
the work reviewed in this chapter. Special thanks to A. Paul Alivisatos, Andrew
Guzelian and Yunwei Cao. [This work was supported by the ISF, BSF and the DIP
programs.]
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42 Miller, J. E., Kidd, K. B., Cowley,
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43 Healy, M. D., Laibinis, P. E., Stupik,
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44 Robinson, W. T., Wilkins, C. J.,
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45 Wells, R. L., Aubuchon, S. R., Sher,
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46 Jin, S., McKee, V., Nieuwenhuyzen,
M., Robinson, W. T., Wilkens, C. J.
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47 Inoue, K., Yoshizuka, K.,
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48 Sato, T., Ruch, R. Stabilization of
Colloidal Despersions by Polymer
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49 Kimura, K. J. Phys. Chem. 1994, 98,
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50 Hamaker, H. C. Recl. Trav. Chim.
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3.1.3
Synthesis and Characterization of Ib-VI Nanoclusters
Stefanie Dehnen, Andreas Eichhöfer, John F. Corrigan, and
Dieter Fenske
Introduction
The synthesis of chalcogen-bridged molecular clusters of the coinage metal elements, copper, silver, and gold, represents an area of increasing activity in recent
chemical and materials science research. This can be attributed in large part to two
facts. Firstly, binary coinage metal chalcogenides feature relatively high ionic [1]
and even higher electric [2] conductivity in the solid state, leading to properties
that are positioned between those of semi-conducting and metallic phases. Secondly, an enormous general interest exists in the size dependency of the chemical,
physical, and structural properties of substances when going from small molecules
to bulk materials [3]. Therefore, numerous research groups are concentrating on
the syntheses and properties of colloidal nano particles with narrow size distributions [4]. An alternative approach is the formation and isolation of crystalline
cluster compounds that are suitable for single-crystal X-ray diffraction. In both
cases, the molecules have to be kinetically protected from further condensation reactions to form the thermodynamically favored binary phases. Therefore, a ligand
sphere coordinating to either surface metal centers or chalcogen atoms is necessary, and this in most cases consists uniquely of tertiary phosphine molecules or
phosphines in combination with other organic groups. Herein we focus on ligandstabilized chalcogen-bridged metal nanoclusters.
3.1.3.1
107
108
3 Syntheses and Characterizations
Experimental [5–33] and quantum chemical [34–37] investigations have been
undertaken during the last decade in order to study the transition from the obviously non-metal Cu2 E monomer (E ¼ S, Se, Te) to bulk Cu2 E [38–40]. From the
synthetic point of view, the merit of the endeavors was the perception of the close
relationship of the chosen reaction conditions with the formation of the desired
products, as well as the use of the acquired knowledge for the synthesis of different
cluster types. Furthermore, the great number of isolated and structurally characterized cluster compounds and comprehensive computational work on this subject
allowed the development and explanation of size-dependent structure principles
observed in the molecules. Lastly, various calculations and measurements have
been carried out in order to get an insight into the physical properties of the cluster
materials [10b–12].
3.1.3.2
Chalcogen-Bridged Copper Clusters
Synthesis Routes
The compounds are – with some exceptions – highly air and moisture sensitive.
Therefore, all reaction steps have to be carried out with rigorous exclusion of air
and moisture in a high purity inert atmosphere by employing standard Schlenkline techniques. All solvents are dried and freshly distilled prior to use. Most
chalcogen-bridged copper compounds have been synthesized according to the
following general reaction pathway and possess one of the formulas given in
Scheme 3-1.
From Scheme 3-1 one recognizes the importance of the CuX to phosphine ratio,
the counterion X of the copper salt, and especially the nature of the phosphine
used. In the first step, the latter complexes to the copper center to result in a coordination complex whose reactivity can be subtly modified by the choice of X and
PR2 R 0 . Furthermore, the driving force of the reactions is the formation of thermodynamically stable SiMe3 X and is thus determined by the choice of X via the bond
+ PR 2R’
CuX
X
R, R’
dppR
E
or
+ dppR
organic solvent
+ E(SiMe3)2
or
[CuX(PR2R’)x]y
or
[CuX(dppR)x ]y
+ RESiMe3
– SiMe 3X
– byproducts
= Cl, OAc (Ac = OCCH3)
= organic group
= bis-(diphenylphosphino)-alkane, -alkene, alkyne or -aryl
= S, Se, Te
main byproducts are Cu2E, E=PR2R’, (Me3Si)2, (RE)2
General synthesis route for the formation of
phosphine-ligated, chalcogen-bridged copper clusters.
Scheme 3-1
[Cu2nEn(PR2R’)m]
or
[Cu2nEn(dppR)m]
or
[Cu2nEn–a(ER)2a(PR2R’)m]
or
[Cu2nEn–a(ER)2a(dppR)m ]
3.1 Semiconductor Nanoparticles
strength of the SiaX bond [41]. Finally, the steric demand of the PR2 R 0 molecules
coating the cluster surface – as measured by Tolman’s cone angle y [42] – is,
though not linearly, closely related to the final size and shape of the copper chalcogenide core.
In addition to these variables, other reaction conditions, such as the chosen solvent and the reaction temperature or temperature program, strongly influence the
course of the synthesis. It has been shown that under similar reaction conditions
the use of derivatives RE(SiMe3 ) (R ¼ organic group) leads to the formation of
completely different structures compared to reactions employing Se(SiMe3 )2 . The
reason for this is the different reactivities of the compared substances as well as a
partly incomplete SeaC bond cleavage. Thus, some RE fragments are bonded at
the periphery of the cluster framework and also act as surface ligands in conjunction with the phosphines.
The copper atoms in the vast majority of the clusters can be assigned a formal
charge of þ1, while the chalcogen ligands are formally viewed as E 2 or RE
groups. Some of the selenium-bridged species, however, and nearly all copper telluride clusters form non-stoichiometric compounds that display mixed valence
metal centers in the formal oxidation states 0 and þI or þI and þII. These observations correlate with those made for the binary phases Cu2 S, Cu2x Se, and
Cu2x Te [38–40].
Even though most ligand-covered copper chalcogenide clusters were prepared
using the above described methods, other approaches for the generation of ligated
or naked [Cu2n En ] cluster compounds have been reported previously. In addition to
syntheses in solvent environments [5–31] that led to the formation of, e.g.,
[PhMe3 As]4 [Cu8 {S2 C2 (CN)2 }6 ] [5], (Me4 N)2 [Cu5 (SPh)7 ] [6], (Ph 4 P)4 [Cu12 S8 ] [7], or
[Cu 6 {Se(2,4,6-iPr3 C6 H2 )}6 ] [8], ligand-free particles have been generated and characterized by means of laser ablation techniques in combination with mass spectroscopy.
Sulfur-Bridged Copper Cluster
The reactions yielding the phosphine-coated sulfur-bridged copper clusters that
have been isolated and structurally characterized to date are summarized in
Scheme 3-2 in order of increasing cluster size.
Sulfur-bridged clusters could not be obtained by reactions starting from copper
halides. The reason is the lower reactivity of a primary copper acetate-phosphine
complex when compared to analogous halide complexes due to chelating acetate
ligands that leave in a double-step substitution of sulfur ligands for oxygen [43].
Reactants of higher reactivity lead directly to the formation and precipitation of
Cu2 S even at low temperatures. The kinetic restrictions at the beginning of the
synthesis combine with those at the end of the cluster growth. The phosphine
molecules have to display a certain minimum steric demand, as it has not yet
proved possible to generate sulfur-bridged copper clusters stabilized with PMe3 ligands using this strategy.
Half the number of known copper sulfide clusters adhere to the general formula [Cu12 S6 (PR2 R 0 )8 ], which represents the smallest known copper chalcogenide
109
110
3 Syntheses and Characterizations
R = R’ = Et
R = Et, R’ = Ph
R = Ph, R’ = Et
Et2O (6: THF)
–80 °C
R = R’ = nPr
R = Ph, R’ = nPr
+ S(SiMe 3)2
CuOAc + PR2R’
warming up to
–25 °C (1, 5-10)
or to room temp.
(2, 3)
– SiMe 3OAc
R = R’ = Ph
R = tBu, R’ = nBu
R = iPr, R’ = Me
R = tBu, R’ = Me
R = tBu, R’ = Me
[Cu12S6(PEt3)8]
1 [13]
[Cu12S6(PEt2Ph)8]
2 [29]
[Cu12S6(PEtPh2)8]
3 [13]
[Cu12S6(PnPr3)8]
4 [15]
[Cu12S6(PPh2nPr)8]
5 [29]
[Cu20S10(PPh3)8]
6 [13]
[Cu20S10(PnButBu2)8] 7 [15]
[Cu24S12(PMeiPr2)12]
8 [17]
[Cu28S14(PtBu2Me)12] 9 [17]
[Cu50S25(PtBu2Me)16] 10 [17]
Survey of the synthesis of sulfur-bridged copper
clusters protected by terminal phosphine ligands.
Scheme 3-2
[Cu2n En ] cluster core (1–5) [13, 15, 18]. The respective molecular structure corresponds to one of two possible isomers that can both be derived from two highly
symmetric polyhedra of metal or chalcogen atoms. The highly symmetric conformation is the reason for the strong preference for ‘‘Cu12’’ species, where an S6 octahedron is penetrated by a Cu12 cubooctahedron. The latter is distorted, however,
because of the coordination of the metal centers to phosphine ligands, as there is
only enough space for eight instead of twelve phosphine molecules at the cluster
surface within reasonable CuaP bond distances. Different distributions of the terminal ligands to the cubooctahedron atoms lead to the formation of isomers I or
II, respectively. Both cluster structures as well as a computed hypothetical ‘‘naked’’,
undistorted [Cu12 S6 ] [35] core are shown in Figure 3-28.
In the ‘‘Cu12’’ clusters, two coordination modes for copper atoms can be observed. The copper atoms that do not bind to phosphine ligands are nearly linearly
coordinated by sulfur neighbors and are positioned either in the equatorial Cu 4
ring of the cubooctahedron (I) or two each in the top and bottom of the Cu 4 planes
(II). Copper atoms that are attached to PR2 R 0 groups are surrounded in a nearly
trigonal planar manner by two sulfur neighbors and one phosphorus atom. The
phosphine molecules formally ‘‘pull’’ the respective metal atoms out of the based
Cu12 polyhedron, whereas the ‘‘naked’’ copper atoms are slightly pushed towards
the cluster center. All sulfur ligands act as m4 bridges between copper atoms. In
molecules of type I, the eight phosphorus atoms are arranged around the cluster
surface in a cubic manner, whereas one finds a P8 delta-dodecahedron in the other
isomer. Both arrangements are counted among the most stable polyhedra of eight
points on a spherical surface.
3.1 Semiconductor Nanoparticles
111
Cu, Ag
S, Se, Te
SeR, TeR (R =organic group)
P
2
4
C
[Cu12S6]
(a)
(b)
Fig. 3-28 Molecular structures of [Cu12 S6 (PEt2 Ph)8 ] (2) and
[Cu12 S6 (PnPr3 )8 ] (4) (without organic groups) and of a
calculated ‘‘naked’’ [Cu12 S6 ] core. (a) General labelling of the
molecular structures in chapter 3.1.3.1.
Quantum chemical investigations with the program system TURBOMOLE [44,
45], employing MP2 [46] and DFT [47, 48] methods, were carried out assigning the
highest possible symmetry in order to explain the obvious experimentally observed
preference for type I clusters in comparison to those of type II [36]. For model ligands PH3 (MP2 and DFT, highest possible symmetry: D4h for type I or S4 for type
II), both structures that represent local minima on the energy hyperface were isoenergetic within the limits of the methods used. On calculating the clusters
[Cu12 S6 (PR3 )8 ] with R ¼ Et or R ¼ Pr according to both structural isomers (DFT,
highest possible symmetry: C4h or S4 ), i.e. one real (1, 4) and one hypothetical
species each, one does not find a direct correlation of a certain type of substituent
at the phosphorus atoms with purely thermodynamic stabilization. Thus, further
kinetic effects such as solubility and mobility of the hydrocarbyl groups at the experimental temperatures (g 0 K) were found to play key roles in the discrimination
of isomers. The same cone angle values of PEt3 and PnPr3 also serve to underline
this conclusion.
A second pair of isomers is formed in the presence of sterically more demanding phosphines, PPh3 or PnButBu2 , that contain 20 copper atoms in the cluster
framework [13, 15]. The positions of eight CuaP groups again determine the shape
of the entire cluster, whose undistorted base can be viewed as a formal condensation product of the two [Cu12 S6 ] units given in Figure 3-28.
A continuation of cluster growth also takes place with PMeiPr2. Being too bulky
for an arrangement of eight molecules around a [Cu12 S6 ] core and not generating a
suitable ligand sphere with six terminal ligands, the phosphine groups force a for-
112
3 Syntheses and Characterizations
2 “[Cu12S6(PMeiPr2)8]”
“– 2 PMeiPr2”
8
Formal dimerization/condensation reaction of two
hypothetical clusters ‘‘[Cu12 S6 (PMeiPr2 )8 ]’’, resulting in the
molecular structure of [Cu24 S12 (PMeiPr2 )12 ] (8) (without
organic groups). Phosphine ligands that are formally lost
during the dimerization/condensation are marked with a cross.
Fig. 3-29
mal dimerization of ‘‘Cu12’’ clusters to the [Cu24 S12 ] core in 8 [17]. Figure 3-29
shows the formal dimerization of two [Cu12 S6 (PR2 R 0 )8 ] clusters of type II, which
results in the molecular structure of 8 after the loss of two phosphine groups.
In contrast to the formal condensation, 2 ‘‘Cu12’’ ! ‘‘Cu20’’ 4 Cu, that leads
to the sharing of a Cu 4 rectangular face in 6, one observes a conjunction via two
Cu3 triangular faces of the two based cubooctahedra. Thus a Cu 6 octahedron is
formed around the inversion center of the cluster. These six copper atoms show a
novel coordination pattern: they are surrounded in a trigonal planar manner exclusively by sulfur neighbors. The four sulfur atoms in the cluster center are found
3.1 Semiconductor Nanoparticles
9
10
Molecular structures of [Cu28 S14 (PtBu2 Me)12 ] (9) and
[Cu50 S25 (PtBu2 Me)16 ] (10) (without organic groups). The atoms
of topologically equal zones are stressed with bold drawing.
Fig. 3-30
to act as m5 bridges, whereas all sulfur ligands in the previously described clusters
and all other additional sulfur ligands in 8 are m4 bridging. Naturally, the m5 bridges
exhibit longer average CuaS bond distances when compared to m4 bridging sulfur
ligands by binding to four copper neighbors within normal distances (2.159(3)–
2.378(3) Å) and possessing one significantly longer CuaS bond each to copper
atoms of the opposite asymmetric unit (2.485(3), 2.509(3) Å).
In the molecular structure of 8, one observes the beginning of the transition
from the so-called ‘‘small’’ copper sulfide clusters to the ‘‘middle-sized’’ species.
Both cluster sizes possess spherical CuaSaP frameworks with copper centers near
the Sn polyhedral edges, but the ‘‘middle-sized’’ examples show some new structural features: coordination numbers other than four for sulfur atoms and copper
centers that bind to more than two sulfur neighbors.
The same observations are made for clusters 9 and 10, and to a greater extent for
the latter [17]. Unlike the clusters described above, clusters 9 and 10 are topologically independent from smaller complexes. They are however themselves structurally related and are indeed produced side by side from the same reaction. The
molecular structures of 9 and 10 are given in Figure 3-30 emphasizing their structural relationship. In Figure 3-31, the based S14 or S25 polyhedra are contrasted.
Besides various bridging modes of the sulfur atoms (m3 –m5 in 9, m4 –m6 in 10)
and the higher numbers of three sulfur ligands around several copper atoms in 10,
113
114
3 Syntheses and Characterizations
Fig. 3-31 S14 (top) and S25 (bottom) polyhedra in 9 and 10,
respectively. The drawn SaS contacts do not represent binding
interactions, but only serve to explain the geometries of the
polyhedra.
both clusters show another new structural feature characterizing ‘‘middle-sized’’
CuaS clusters. Only for these two largest examples of copper sulfide clusters does
one observe sulfur atoms inside an outer CuaS shell, generating inner CuaS units.
In 9, one sulfur ligand centers a distorted S13 deltahedron with four copper atoms
bonded to the central sulfur atom to yield an inner [SCu 4 ] unit. The remaining Cu
centers are positioned either slightly below the edges of the S13 polyhedron if they
show no CuaP bonds or slightly above if bonded to phosphine ligands. Near C3
symmetry of the CuaS core is distorted by an ‘‘irregular’’ position of a single copper atom that also promotes the m3 or m5 bridging modes of two sulfur ligands. The
spherical cluster core is enclosed by a distorted icosahedron of the twelve phosphorus atoms. The sulfur substructure of 10 represents a cylinder-shaped polyhedron of near D5d symmetry with a length of 17.14 Å. Underlining the supposition that 10 is formed from 9 through rough dimerization-condensation, three
inner sulfur atoms are present. Two of them – formally arising from the central
sulfur atom in 9 – are also m4 bridging, resulting in an [SCu 4 ] fragment. One of
the corresponding copper centers in each of these units, however, belongs to the
cluster surface rather than remaining an inner metal atom because of structural
differences between 9 and 10 when approaching the center in 10. Here one observes the most unusual feature of any of the copper sulfide clusters reported: the
central sulfur ligand binds to six copper atoms in a slightly corrugated chair type
3.1 Semiconductor Nanoparticles
500
400
–1
En [kJ·mol ]
300
200
100
1
10
5
n
Fig. 3-32 Course of the stabilization energies per monomer
unit En in ‘‘naked’’ clusters [Cu2n Sn ] (n ¼ 1–4, 6, 10). The given
values consider the total energies of the most stable structural
isomers computed on MP2 level.
manner. This arrangement, which is unknown for all other copper sulfide compounds, displays large CuaS distances in the range 2.633(4)–2.673(5) Å, which are
about 0.3–0.5 Å longer than usual CuaS bonds. Considering the next ten copper
atoms at distances of within 2.896(5)–3.814(5) Å around the sulfur center, a spherical [SCu16 ] arrangement results in the cluster hub. The structural observations
give evidence for an interstitial S 2 anion rather than a more covalently bonded
sulfur ligand.
By a comprehensive ab initio study of clusters [Cu2n Sn (PH3 )m ] (n ¼ 1–4, 6, 10;
m ¼ 0, 2, 4, 6, 8) with an MP2 level (program system TURBOMOLE), stabilization
energies for the sequential attachment of one Cu2 S unit to a given ‘‘naked’’ cluster
(m ¼ 0, Figure 3-32) as well as the course of the phosphine binding energies for
clusters containing up to twelve copper atoms (m > 0, Figure 3-33) were developed
[35]. For ‘‘naked’’ clusters, the stabilization energies increase continuously with the
cluster size, which proves the investigated molecules to be thermodynamically
unstable species in vacuo, and thus have to be stabilized by ligands. Moreover, the
increment of the stabilization energy per monomer unit decreases with increasing
cluster size as expected, since it should theoretically converge with infinite cluster
size, i.e. solid Cu2 S.
As a third result, one gathers from the calculations that the CuaP binding energy per CuaP bond decreases rapidly from a value of ca. 140 kJmol1 for
[Cu2 S(PH3 )2 ] – with a short relaxation when the Cu:P ratio is decreased (1 for
n ¼ 1–3, >1 for n ¼ 4, 6) – to a value of ca. 56 kJmol1 for the ‘‘Cu12’’ species.
Even if one assigns a higher CuaP binding energy to the ‘‘real’’ tertiary phosphines
with organic side chains, it is evident that the phosphine ligands are bonded very
115
116
3 Syntheses and Characterizations
500
400
300
–1
Eb [kJ·mol ]
200
100
0
1
2
3
4
5
6
n
Fig. 3-33 Course of the CuaP binding energies
per CuaP bond Eb in phosphine-ligated clusters
[Cu2n Sn (PH3 )m ] (n ¼ 1–4, 6; m ¼ 2, 4, 6, 8).
The given values consider the total energies of
the most stable structural isomers computed
on MP2 level with respect to the total energy of
PH3 calculated by the same method.
weakly to the copper sulfide core. Therefore, an additional stabilization by solvents
or a crystal lattice is required.
Selenium-Bridged Copper Clusters
Scheme 3-3 summarizes the reactions yielding all PR2 R 0 -coated selenium-bridged
copper clusters that have been isolated and structurally characterized to date, in
order of increasing cluster size.
In contrast to the sulfur-bridged compounds, the synthesis of selenium-bridged
clusters can start from either copper acetate or copper chloride. This can be ascribed to the lower reactivity of Se(SiMe3 )2 when compared to the sulfur analog
[49, 50]. Thus, the formation of Cu2 Se does not occur as rapidly as the formation
of Cu2 S, and the cluster growth and ligand protection can often be observed even
at room temperatures. However, by reacting copper acetate and by choosing low
temperatures, one obtains smaller clusters that can be characterized as being intermediates during the course of the formation of larger molecules. In addition to
a great number of ‘‘stoichiometric’’ compounds that formally contain Cuþ and
Se 2 , some clusters were characterized that exhibit a Cu:Se ratio less than two, indicating mixed valence metal centers. All selenium ligands feature the formal
charge 2, as one does not find any SeaSe binding interactions and one can
therefore assign Cu I and Cu II within the same cluster core in 14, 17, 27, and 33.
The phosphine ligands again play a key role in the determination of the observed
cluster size and shape; however, the number of known compositions is approximately four times that of the copper sulfide system and the cluster growth pro-
3.1 Semiconductor Nanoparticles
R = Ph, R’ = Et
R = R’ = nPr
R = R’ = nPr
R = R’ = Et
R = Et, R’ = Ph
R = R’ = Cy
R = R’ = iPr
R = R’ = iPr
Et2O or THF
CuX + PR2R’
(X = Cl, OAc)
+ Se(SiMe 3)2
R = tBu, R’ = Me
R = tBu, R’ = Me
R = R’ = Ph
– SiMe 3X
R = R’ = tBu
R = Et, R’ = Ph
R = tBu, R’ = nBu
R = Me, R’ = Ph
R = R’ = Ph
R = R’ = Ph
R = iPr, R’ = Me
R = tBu, R’ = Me
R = R’ = Et
R = Et, R’ = Ph
R = R’ = Ph
R = R’ = Cy
R = R’= Et
R = Et, R’ = Ph
R = R’ = Et
R = R’= Ph
[Cu12Se6(PEtPh2)8]
11 [13]
[Cu12Se6(PnPr3)8]
12 [29]
[Cu12Se6(PCy3)6]
13 [18]
[Cu20Se13(PEt3)12]
14 [10b]
[Cu26Se13(PEt2Ph)14]
15 [18]
[Cu26Se13(PCy3)10]
16 [18]
[Cu29Se15(PiPr)12]
17 [10a]
[Cu30Se15(PR2R’)12]
[Cu31Se15(SeSiMe3)(PtBu2Me)12]
19 [17]
20 [17]
[Cu32Se16(PPh3)12]
21 [21]
[Cu36Se18(PtBu3)12]
22 [10a]
[Cu44Se22(PEt2Ph)18]
23 [14]
[Cu44Se22(PnButBu2)12]
24 [14]
[Cu48Se24(PMe2Ph)20]
25 [17]
[Cu52Se26(PPh3)16]
26 [21]
[Cu59Se30(PCy3)15]
27 [18]
[Cu70Se35(PR2R’)21]
28 [29]
29 [17]
[Cu70Se35(PEt3)22]
30 [10b]
[Cu70Se35(PEt2Ph)n](n = 23, 24) 31 [26]
[Cu72Se36(PPh3)20]
[Cu74Se38(PCy3)18]
[Cu140Se70(PR2R’)34]
32 [21]
33 [31]
34 [23]
35 [31]
[Cu140Se70(PEt3)36]
36 [23]
[Cu146Se73(PPh3)30]
37 [12]
Survey of the synthesis of selenium-bridged
copper clusters protected by terminal phosphine ligands.
Scheme 3-3
18 [10a]
117
118
3 Syntheses and Characterizations
ceeds to a much higher extent. Again, the lower tendency to precipitate the binary
phase can be held responsible for this observation.
Another route into selenium-bridged copper polynuclear complexes or clusters is
their synthesis using RSeSiMe3 (R ¼ organic group) in the presence of either tertiary phosphines or chelating ligands such as 1,2-bipyridyl, diphenylphosphino
acid or bidentate phosphines dppR (R ¼ organic spacer). The compounds thus
prepared are listed in Scheme 3-4 and ordered according to their core size.
Again, one finds clusters containing exclusively Cu I , as well as compounds with
mixed-valence metal centers. In contrast to the compounds presented in Schemes
2 and 3, one example of a selenido/selenolato-bridged cluster exists apparently
containing Cu I and Cu 0 . Apart from this peculiarity, the cluster size and conformation of the Se 2 /SeR -bridged compounds is now not only influenced by the
nature and steric demand of the phosphine used, but it is also dependent on the
spatial properties of the organic substituent at the selenium reactant. This determines both the Se 2 :SeR ratio and the ensuing arrangement of Se 2 ligands
positioned inside the cluster core and SeR ligands at the periphery of the molecules.
The delicate, non-linear structural influence of the phosphine ligands allows for
the formation of different clusters of often significantly varying size and shape in
the presence of the same phosphine. The isolation of various cluster compounds is
then enabled only by the judicious choice of reaction conditions. In contrast, the
use of different PR2 R 0 ligands can lead to clusters of the same [Cu2n Sen ] cores that
possibly display differing numbers of terminal ligands and/or structural isomerism. These observations have already been illustrated and explained for the sulfurbridged copper clusters, but they are more pronounced for the CuaSe species, and
a much higher structural variety is found. As for the copper sulfide compounds,
one can discriminate between ‘‘small’’ (11–13) and ‘‘middle-sized’’ (14–27) clusters
with respect to their shape and the coordination properties of the atoms concerned. However a novel group of ‘‘large’’ clusters (28–37) exists for the CuaSe
system. Whereas the transition from the first to the second type of molecular size
occurs almost seamlessly, a clear break occurs when the number of copper atoms
exceeds 59.
The only composition and structure that is identical for both the copper sulfide
and selenide clusters is that for [Cu12 E6 (PR2 R 0 )8 ] (E ¼ S, Se; PR2 R 0 ¼ PEtPh2 ,
PnPr3; respective cone angles: 141, 132 ). Compounds 11 [13] and 12 [29] are topologically identical to their sulfur analogs. The use of a much larger phosphine
ligand PCy3 does not lead to a condensation or dimerization of cluster cores as
observed for 6–8 with PR2 R 0 groups of medium or large size around. Unlike the
copper sulfide system, in the presence of PPh3 , PnButBu2 , or PMeiPr2 a reduction
in the number of terminal ligands to six is acceptable in the case of 13 [18], and the
resulting ligand shell is suitable for the cluster core. The clearly larger cone angle
of PCy3 when compared to PPh3 , PnButBu2 , or PMeiPr2 (170 vs 145, 146, and
165 ) allows a distorted P6 octahedron to sufficiently protect the copper selenide
framework, although a third variation of a distortion of the [Cu12 E6 ] core is observed.
3.1 Semiconductor Nanoparticles
R = bd
R=m
119
[Cu2(SePh)2(PPh3)3]
38 [30]
[Cu2(SePh)2(dppbd)3]
39 [20]
[Cu3(SeMes)3(dppm)]
40 [22]
[Cu6(SePh)6(bipy)2]
41 [30]
THF or toluene
[Cu9(SePh)6(O2PPh2)3]
42 [30]
+ PhSeSiMe3
R = be
or MesSeSiMe 3 (40)
43 [20]
[Cu16Se4(SePh)8(dppbe)4]
or nBuSeSiMe3 (45)
R
=
p
or tBuSeSiMe3 (51)
[Cu(dppp)2][Cu25Se4(SePh)18(dppp)2]
+ additional
Se(SiMe3)2 (53)
[Cu32Se7(SenBu)18(PiPr3)6] 45 [23]
CuX + dppR
– SiMe3X
or PPh3 (38)
R=a
or PiPr3 (45, 51, 53)
46 [22]
[Cu36Se5(SePh)26(dppa)4]
or bipy (41)
or Ph2PO2H (42)
R=e
[Cu36Se13(SePh)12(dppR)6] 47 [30]
R=b
48 [30]
49 [30]
R=p
[Cu38Se13(SePh)12(dppR)6]
50 [22]
R=b
[Cu50Se20(SetBu)10(PiPr3)10] 51 [23]
R=a
Mes = C6H3Me
X
= Cl, OAc
dppR = µ-Ph2P-R-PPh2
= bis-(diphenylphosphino)-alkane
-alkene
-alkyne
-aryl
organic spacers R:
m= methane
e = ethane
p = propane
b = butane
a = acetylene
bd = 1,3-butadiene
be = benzene
Scheme 3-4 Survey of the synthesis of selenido/selenolatobridged copper clusters protected by monodentate or chelating
bidentate ligands.
[Cu58Se16(SePh)24(dppa)6]
52 [22]
[Cu73Se35(SePh)3(PiPr3)21]
53 [23]
44 [22]
120
3 Syntheses and Characterizations
14
15
16
Molecular structures of [Cu20 Se13 (PEt3 )12 ] (14),
[Cu26 Se13 (PEt2 Ph)14 ] (15), and [Cu26 Se13 (PCy3 )10 ] (16) (without
organic groups).
Fig. 3-34
With the clusters discussed, one perceives the strong preference for highly symmetric substructures of all components involved in the heavy atom framework.
Capped or uncapped, sometimes condensed octahedral, prismatic, cubic, dodecahedral, cubooctahedral, or icosahedral arrangements of either metal, chalcogen or
phosphorus atoms penetrate or enclose each other in all structures. Among the
discussed compounds, the ‘‘Cu12’’ molecules possess the most symmetric substructures. Selenium ligands, however, evidently prefer higher bridging modes.
Therefore the formation of ‘‘small’’ clusters with a strong m4 bridging restriction is
disadvantaged and thus rarer than the formation of larger ones. For the CuaSe
clusters, three additional ‘‘middle-sized’’ compounds are interesting with respect to
their symmetry properties as well as being an additional illustration of the phosphines’ structural influence. The cluster frameworks of 14, [10b] and the CuaSe
isomers 15 [18], and 16 [18] can again be described as an alternating packing of
copper, selenium and phosphorus polyhedra, respectively. Regular polyhedra are,
however, only observed in the structures of 14 and 15, whereas the [Cu26 Se13 ] isomeric core of 16, being surrounded by only ten PCy3 (versus 14 PEt2 Ph groups),
does not follow a known polyhedral pattern. The molecular structures of 14–16 are
shown in Figure 3-34.
Clusters 14 and 15 are topologically related, as both contain a centered icosahedron of selenium atoms, although the penetrating polyhedra of copper atoms are
arranged in different manners according to the different number of copper atoms
and CuaPR2 R 0 fragments (14: R ¼ R 0 ¼ Et, 15: R ¼ Et, R 0 ¼ Ph) in both clusters
(Figure 3-35).
3.1 Semiconductor Nanoparticles
Fig. 3-35 Comparison of the penetrative polyhedra in 14 (top)
and 15 (bottom) that form the heavy atom cluster cores. The
drawn lines only represent polyhedral edges and do not specify
binding interactions.
In 14, a Cu8 cube surrounds the inner selenium atom below the Se3 faces of
the Se12 icosahedron. An outer Cu12 icosahedron follows, which is positioned
above and perpendicular to the Se12 shell. The eight copper centers inside the
cluster thus form tetrahedral [CuSe4 ] groups, and the twelve outer ones display a
[Se3 CuP] environment by binding to the equally orientated P12 icosahedron of the
PEt3 ligands. In 15, the inner selenium atom has twelve copper next nearest
neighbors that form an icosahedron, each edge being positioned below an SeaSe
edge of the enclosing Se12 polyhedron and perpendicular to it, resulting in an
coordinative occupation of one copper atom per Se3 face. Slightly above the Se12
icosahedral surface, an arrangement of 14 CuaPEt2 Ph units can be described as a
sixfold capped cube. The twelve inner copper atoms and the eight metal centers of
the capped cube that are all placed near Se3 faces are approximately tetrahedrally
coordinated by either four selenium ligands or three selenium atoms and one
phosphorus ligand – characteristic of clusters being larger than ‘‘small’’ ones. The
six remaining CuaPCy3 fragments are situated above six SeaSe edges and therefore represent three-coordinated [Se 2 CuP] units, common for both ‘‘small’’ and
121
122
3 Syntheses and Characterizations
‘‘middle-sized’’ clusters. When comparing the Cu:Se ratio for 14 and 15, one
notices the fundamental difference that may be responsible for the structural differences. Cluster 14 represents a mixed-valence compound featuring formally six
Cu II and fourteen Cu I centers to balance the 26 charge of the Se13 substructure.
The intermetallic distances in 14 however (Cu(1)aCu(1): 2.821(4), Cu(1)aCu(2):
2.678(4) Å) do not suggest significant interatomic interactions despite the partial
deficiency of electrons. The only structural hint of different valence situations in 14
or 15 is that, apart from equal average CuaP bond lengths (2.247 Å in 14 and 15)
and similar average intermetallic distances CuaCu (2.707 Å in 14, 2.682 Å in 15),
one measures significantly shorter average CuaSe bonds in the mixed-valence
compound 14 (2.365 Å) versus 15 (2.457 Å).
The molecular structure of 16 differs from those of 14 and 15, since a very
irregular Se13 deltahedron forms the selenium substructure. Twelve of the copper
centers and ten CuaPCy3 fragments are bonded to the Se13 polyhedron in such a
way that they are again three- or four-coordinate. Unlike the situation in 14 or 15,
one also observes four copper atoms that show a nearly linear coordination to two
selenium neighbors. This pattern is typical of ‘‘small’’ CuaE clusters and again
underlines the partial propagation of these properties in the ‘‘middle-sized’’ compounds.
Similar to the ‘‘Cu12’’a‘‘Cu20’’, ‘‘Cu12’’a‘‘Cu24’’ and ‘‘Cu28’’a‘‘Cu50’’ relationships
discussed for copper sulfide clusters, the molecular structures of 14 and 15 prove
to be closely related, even though the numbers and ratios of heavy atoms are very
different. In contrast, a similar number of copper and selenium atoms can, but
need not necessarily lead to a topological relationship. Compounds 17–21 [10a, 17,
21] contain 29, 30, 31 and 32 copper atoms, respectively, in addition to 15 or 16
selenium ligands. Nevertheless, three completely different molecular frameworks
are observed. The first structural type is adopted by 17 [10a] , 18 [10a] , and 19 [17].
As an example, the molecular structure of 17 is given in Figure 3-36.
17
Fig. 3-36
Molecular structure of [Cu29 Se15 (PiPr)12 ] (17) (without organic groups).
3.1 Semiconductor Nanoparticles
20
Fig. 3-37 Molecular structure of [Cu31 Se15 (SeSiMe3 )
(PtBu2 Me)12 ] (20) (without organic groups).
Although the number of copper atoms is reduced by one in 17 in comparison to
18 or 19, indicating a mixed valence situation, all clusters show identical CuaSeaP
cores. The additional copper atom present in 18 and 19 is positioned in the cluster
center. A reaction employing the same reactants and temperatures as for the synthesis of 19, but using a 40% higher PtBu2 Me:CuOAc ratio, yields compound 20
[17]. Analogous atomic arrangements or corresponding polyhedra are not found in
19 and 20, even though the cluster core of 20 (Figure 3-37) can equally be described
as consisting of five parallel planes of copper and selenium atoms that are surrounded by twelve phosphine ligands.
The SeSiMe3 group in 20, whose occurrence is unique for copper chalcogenide
clusters, not only influences the molecular structure of 20 but it also clears up the
obvious reason for the structural differences for this series. Cluster 20 can be
viewed as a ‘‘frozen’’ intermediate on the way to the formation of a larger cluster
such as the ‘‘Cu70’’ cluster 29 [17] (see below), as a leaving group is still present.
The latter uniquely acts as a protecting unit with the phosphine ligands, and
therefore further cluster growth is slowed down to such an extent that the compound crystallizes preferentially. Its formation is thus naturally independent of the
cluster growth yielding 19. A third structural type for similar cluster sizes is found
in 21 [21], in which the selenium atoms form a flattened polyhedron consisting of
non-bonded Se3 triangles. The copper atoms show quasi-linear, distorted trigonal
planar or tetrahedral coordination geometries. These coordination patterns and m5
and m6 bridging modes of the selenium ligands are common for ‘‘middle-sized’’
CuaE clusters. Complexes 17 and 21 represent exceptions for their cluster size in
another way, as they do not contain copper or selenium atoms inside a spherically
tolerable CuaSe shell. All other ‘‘middle-sized’’ clusters with 20 copper atoms and
123
124
3 Syntheses and Characterizations
Selenium substructures in 22–27 (from top left to
bottom right). The drawn SeaSe contacts do not represent
bonding interactions, but rather illustrate the geometries of
the polyhedra.
Fig. 3-38
upwards (beginning with 14) have either central copper atoms (18, 19) or central
selenium ligands (14–16, 20).
The larger molecules 22–27 [10a, 14, 17, 21, 18] also enclose selenium atoms
within a Sen polyhedron (n ¼ 18, 22, 24, 26, or 30). The number of inner selenium
ligands consequently increases with increasing cluster size from one central atom
(22) [10a] to two inner ones (23–25) [14, 17] to three selenium centers inside the
cluster core (26, 27) [21, 18]. The respective underlying selenium substructures are
summarized and contrasted in Figure 3-38.
Although these six compounds are all classified as being spherical ‘‘middlesized’’ clusters, and all have central units within a CuaSe surface, one can distinguish two different shapes. Clusters 22 and 26 (Figure 3-39) adopt a triangular
topology, whereas 23–25 and 27 form ellipsoid-type cluster surfaces (Figures 13
and 14).
The molecular structures of both 22 and 26 display idealized C3 symmetry that
is not realized in the crystal (22: C1 symmetry, 26: C2 symmetry). Along the idealized threefold axis one finds a selenium ligand at the top and at the bottom of the
flattened trigonal prism, as roughly formed by the cluster atoms. In 22, a third
selenium ligand is positioned between the latter two in the cluster center, whereas
in 26, three inner selenium atoms are grouped around the center of the CuaSe
core and two copper atoms are arranged along the pseudo C3 axis between the two
above-mentioned selenium ligands. Beginning with the central unit, the cluster
is formed by the usual alternating bonds from copper to selenium atoms. For 22,
[Se 2 Cu], [Se3 Cu], or [Se 2 CuP] coordination environments around copper centers
and m3 –m6 bridging selenium ligands result, with CuaSe distances of between
2.347(4) and 2.875(5) Å. The larger cluster core of 26 features consequently higher
3.1 Semiconductor Nanoparticles
26
22
Molecular structures of [Cu36 Se18 (PtBu3 )12 ] (22) and
[Cu52 Se26 (PPh3 )16 ] (26) (without organic groups).
Fig. 3-39
23
24
Molecular structures of [Cu44 Se22 (PEt2 Ph)18 ] (23)
and [Cu44 Se22 (PnButBu2 )12 ] (24) (without organic groups).
Fig. 3-40
coordination numbers that show near-tetrahedral [Se3 CuP] arrangements in addition to two- and three-coordinated copper atoms as in 22, and the selenium atoms
act as m4 –m8 bridges within a wide CuaSe range of distances ranging from 2.19 to
2.91 Å. Twelve PtBu3 ligands enclose the copper selenide core of 22, and sixteen
PPh3 ligands are bonded to the CuaSe core of 26.
The structural isomers 23 and 24 are obtained via completely different reaction conditions including the use of phosphine ligands, different molar ratios of
CuOAc:PR2 R 0 , and the exposure of the reaction mixtures to different final reaction
temperatures.
125
126
3 Syntheses and Characterizations
25
27
Molecular structures of [Cu48 Se24 (PMe2 Ph)20 ] (25)
and [Cu59 Se30 (PCy3 )15 ] (27) (without organic groups).
Fig. 3-41
The spatial demands of the phosphine ligands allowed ligand shells of either
eighteen PEt2 Ph or twelve PnButBu2 groups to sufficiently stabilize a [Cu44 Se22 ]
core. In 23, the phosphorus atoms form a regular deltahedron of alternating,
eclipsed rings P3 , P6 , P6 , and P3 along the long axis of the cluster. The twelve
phosphorus atoms at the periphery of the cluster core in 24 define a highly distorted icosahedron. As a result of the different number and thus arrangement of
PR2 R 0 ligands that expectedly affect the core geometry by ‘‘pulling’’ the attached
copper atoms somewhat out of the CuaSe core, some differences are observed on
examining the copper selenide frameworks in detail. This is clearly seen when the
respective Se22 substructures are compared with each other (Figure 3-38). Apart
from different distortions of the described peripheral Se20 polyhedron (total
lengths or widths of the Se22 substructures according to the orientations in Figure
3-41: 11.43 8.64 Å in 23, 8.34 10.60 Å in 24), the relative positions of the inner
Se 2 ‘‘dumbbells’’ are perpendicular to each other. One can therefore assign a fragment of twelve copper and two selenium atoms in both cluster cores that are positioned completely differently, which is initiated by the significantly different situations on the cluster surface mentioned above. The isomerism observed with the
‘‘Cu44’’ clusters again shows that certain [Cu2n En ] compositions represent ‘‘isles of
stability’’ – within given limits of structural isomers – that are experimentally obtained whenever a suitable ligand shell can be realized by means of the selected
phosphine molecules.
Two further ellipsoid-like arrangements are found for clusters 25 and 27 (Figure
3-4). Apparently, it is not possible to create a phosphine shell for surrounding a
[Cu44 Se22 ] cluster by a suitable number and arrangement of PMe 2 Ph groups (cone
angle: 127 ). Instead, the smallest compound to be stable and insoluble enough for
isolation is the slightly larger cluster 25, which is however structurally related to
3.1 Semiconductor Nanoparticles
the somewhat smaller species 23. The ellipsoid-like Se24 substructure of 25 (Figure
3-38) can be derived from the Se22 arrangement in 23 via substitution of two selenium ligands capping the outer Se6 rings for the single Se center in the smaller
polyhedron. Again, one finds all coordination features typical of ‘‘middle-sized’’
CuaSe clusters, with four copper atoms that show a nearly tetrahedral environment of four selenium ligands. Unlike all smaller clusters that contain selenium
atoms inside the cluster core, 25 displays two tetrahedral [SeCu 4 ] units – linked by
a trans-edge junction via a compressed Cu 4 tetrahedron – with a uniquely small
bridging number of four at the two inner selenium ligands (cf. bridging modes of
the inner selenium ligands in 14–16, 20, 22–24: m5 –m8 , m12 ). Together with 14 and
17, 27 is the third copper selenide cluster that displays a number of metal centers
that is too small for an overall þI oxidation state at the copper atoms. A single
Cu 2þ center cannot however be assigned and the ‘‘lack’’ of one electron is likely to
be delocalized over the cluster framework. As in 23–25, the molecule is oblong in
shape and represents the largest spherical copper selenide cluster known. It concludes the series of ‘‘middle-sized’’ CuaSe clusters. The selenium atoms form a
highly irregular Se27 polyhedron around three inner chalcogen centers (Figure 338) that can alternatively be viewed as a rough arrangement of selenium ligands in
three layers. The molecular structure of 27 includes all possible coordination environments around copper atoms, [Se 2 Cu], [Se3 Cu], [Se4 Cu], [Se 2 CuP], and [Se3 CuP]. Selenium ligands are observed to act as m4 , m5 , m6 , or m7 bridges.
Summarizing the structural properties of the ‘‘small’’ CuaSe clusters 11–13 and
the ‘‘middle-sized’’ ones 14–27, one perceives the spherical shape and the absence
of any topological relationship with that of the bulk material Cu2 Se as a central
feature. Only the tendency toward higher coordination numbers three and four at
the copper centers with increasing cluster size approaches the situation in the binary phase. The principal difference between ‘‘small’’ and ‘‘middle-sized’’ clusters
is the occurrence of inner atoms that appears firstly with a copper selenide core of
33 atoms (14). Typical of the coordination patterns in the ‘‘middle-sized’’ molecules
is a mixture of higher and lower coordination numbers or bridging modes, the
lower being observed in ‘‘small’’ clusters without variation.
Although only eleven additional copper atoms are observed in compounds 28–31
[17, 10b, 29, 26], they represent the smallest examples of a new type of copper
selenide clusters. The so-called ‘‘large’’ CuaSe clusters are characterized by significantly different cluster shapes, structural principles and relationships to the
bulk material when compared to the smaller species. Clusters 28–37 display A-B-Atype packed layers of selenium atoms that form large triangular frameworks. A
weak image of the molecular structures of 22 or 26 arises; however, a closer examination clearly enables the latter series to be distinguished from the layer-type
structures. Clusters 28–37 are all based on three such selenium layers, of which
the middle one possesses the most chalcogen atoms. Figure 3-42 shows the selenium networks of the ‘‘large’’ clusters 28 [29] (as an example of the Cu70 Se35
cores), 32 [21] (Cu72 Se36 core), 33 [31] (Cu74 Se38 core), 34 [23] (as an example of
the Cu140 Se70 cores), and 37 [12] (Cu146 Se73 core). It is worth mentioning that this
structural pattern seems to be exceptionally advantaged for large copper selenide
127
128
3 Syntheses and Characterizations
Fig. 3-42 Selenium substructures in 28, 32,
33, 34, and 37 (from top left to bottom right),
which contain 35, 36, 38, 70, or 73 selenium
atoms. The drawn SeaSe contacts do not
represent binding interactions, but only serve
to visualize the A-B-A packing of selenium
atoms in three layers, which is additionally
demonstrated by the supplementary view of
the Se36 substructure of 32 (top mid). Open
circles represent the positions of disordered
copper atoms (see text and following figures).
clusters, as it is observed equally for a size range from ‘‘Cu70’’ to ‘‘Cu146’’, i.e. even
with a doubling of the molecular mass.
Clusters 28–31 contain Se10 , Se15 and Se10 layers that represent perfect triangles.
The next larger selenium substructure should be formed by an arrangement of
Se15 , Se21 , and Se15 triangular layers, each displaying one more row of selenium
atoms. A corresponding ‘‘[Cu102 Se51 ]’’ cluster core is as yet unknown, and one observes instead the formation of smaller clusters in between this range that show
selenium layers with the middle one deviating from a perfect triangular shape. In
32, one corner of the middle layer is modified with respect to the corresponding
one in the ‘‘Cu70’’ clusters in such a way that two selenium atoms replace one. All
three corners are similarly changed, to result in the Se10 aSe18 aSe10 selenium
framework of 33, and this compound is another cluster featuring fewer copper
centers than Cu2n with respect to Sen . The third ‘‘idealized’’ cluster composition is
actually realized with 34, 35 and 36. The underlying selenium layers contain 21,
28, and 21 atoms. By replacing all corner selenium centers of the middle triangle
by two chalcogen atoms, one obtains the Se73 substructure that is observed in 37.
For this third possible size, a successive substitution of triangle corners in the
middle layer could not as yet be experimentally proven, whereas for the first,
smaller size (35 selenium atoms), a modification of one or three corners was
observed.
The reason for the missing ‘‘Cu102’’ cluster might be related to the primary steps
of the cluster growth mechanism. Selenium networks consisting of three triangle
3.1 Semiconductor Nanoparticles
hexagonal layers based on 0.5[n(n þ 1)] selenium atoms (n ¼ number of straight
rows in the triangle network) in A-B-A order offer three possible situations for the
center of the triangle: one either finds a hole (representing the unoccupied C
position), one selenium atom in the middle layer (located on a B position), or
two selenium atoms belonging to the two outer layers (located on A positions).1)
Since one observed only such ‘‘large’’ CuaSe cluster that follow the last two
possibilities, the presence of a central Cux Sey axis (SeA aCuaCuaSeA in 28–33,
CuaSeB aCu in 34–37) seems to be necessary for the formation of a stable cluster
core. Assuming that the cluster growth begins in the center of the molecule, the
presupposition of such a [Cu2 Se] or a [Cu2 Se 2 ] fragment being situated at this
place appears feasible for a non-spherical cluster. Therefore, only those molecules
that contain either one (middle; cf. 34–37) or two (outer; cf. 28–33) selenium
layers with a central selenium atom instead of a central triangular hole are formed
and experimentally observed. Indeed, neither Se15 nor Se21 layers that would be
exclusively underlying a ‘‘Cu102’’ molecule contain central selenium atoms and
should therefore be only found together with other triangles.
In Figure 3-43, the cores of the clusters containing 70 (28–31), 72 (32) and 74
(33) copper atoms are compared.
The ‘‘Cu70’’ core is quite tolerant with regard to its ligand shell. An increase in
the spatial demand of the PR2 R 0 molecules is not strictly coupled with a decreased
number of terminal ligands. However, one still finds a rough correlation, since
twenty-one groups of the bulkier phosphines PMeiPr2 (cone angle: 146 ) and
PtBu2 Me (cone angle: 161 ) are arranged around the CuaSe core, whereas twentytwo and twenty-three terminal ligands are observed for the smaller phosphines
PEt3 or PEt2 Ph (cone angles: 132 or 136 respectively). The clusters 28, 29 [17], and
31 [26] differ marginally with respect to their molecular structures. In 28 and 29,
the only observable difference is a slightly less complete ligand shell achieved
by the less demanding PMeiPr2 groups. Cluster 31 displays a disorder problem
with the PEt2 Ph groups, such that it cannot be determined with certainty whether
the total number of phosphine ligands sums to twenty-three or twenty-four. Accordingly, for the formation of the cluster, another two (or three) additional phosphine molecules – versus 28 or 29 – are formally arranged in the triangle corners
in such a way that two (or three) corners possess four instead of three CuaPEt2 Ph
groups and one corner (or none) retains three of them. An intermediate number of
phosphine ligands, twenty-two, is observed with a more distinct structural modification, but it is again the corner region that is affected. In contrast to all of the
corner arrangements of copper and selenium atoms in the other ‘‘Cu70’’ clusters
that cannot be described by means of regular patterns, in 30 [10b] only two corners
follow this pattern. The third shows a highly ordered [Cu9 Se3 (PEt3 )4 ] configuration featuring a centered cube of copper atoms that is capped at three faces by
m4 -selenium ligands and carries four phosphine groups at the bottom atoms. This
1) Respective layers are potentially one of every third
number 0.5[n(n þ 1)], i.e. Se1 , Se10 , Se28 , etc., whereas
the layers that feature a central hole show selenium
numbers in between, i.e. Se3 , Se6 , Se15 , etc.
129
130
3 Syntheses and Characterizations
28, 29
32
Molecular structures of
[Cu70 Se35 (PR2 R 0 )21 ] (R ¼ iPr, tBu; R 0 ¼ Me)
(28, 29), [Cu70 Se35 (PEt3 )22 ] (30),
[Cu70 Se35 (PEt2 Ph)n ](n ¼ 23, 24) (31),
[Cu72 Se36 (PPh3 )20 ] (32), and [Cu74 Se38 (PCy3 )18 ]
Fig. 3-43
31
30
33
(33) (without organic groups). Two copper
atoms that are positioned around the central
SeaCuaCuaSe axis are statistically distributed
with a 0.67 occupation over three positions in
28, 29 or 31.
peculiarity of one cluster corner in 30 is one of two reasons for the missing
pseudo-threefold molecular symmetry. A C3 symmetry of the [Cu70 Se35 ] core of 28,
29 or 31 is additionally approached – apart from one ‘‘wrong’’ CuaPR2 R 0 position
in 28 or 29 – by the statistical disorder of two copper atoms. The latter are distributed over three positions – with an accordant 0.67 occupancy – around the
CuaSeaSeaCu pseudo-threefold axis running through the center of the molecule.
The mentioned disorder situation of PEt2 Ph ligands in 31 also leads to a pseudo-C3
axis as far as 24 phosphine groups are assigned to the ligand shell. In contrast,
a close look to the cluster center in 30 shows an open copper atom position. The
two other metal atom sites are fully occupied, leading to a lower – but crystallographically realized – C2 symmetry.
The molecular structure of 32 can be derived from that of 28 or 29 in a simple
manner if one firstly fully occupies all three positions of copper atoms around the
CuaSeaSeaCu axis mentioned before, and, secondly, one formally substitutes one
[Se 2 CuPPh3 ] unit for one corner selenium atom. In this way, the modification of
the central selenium layer occurs. The number of phosphine groups that enclose
the respective corner is additionally reduced to two, since two of the three copper
atoms that bind PR2 R 0 ligands in the analogous corners in 28 or 29 remain
‘‘naked’’ in 32. Such (formal) alteration of the cluster corner, and full occupation of
the Cu3 group around the cluster center, cause the formation of the mixed-valence
3.1 Semiconductor Nanoparticles
36
34
37
Molecular structures of
[Cu140 Se70 (PEt3 )34 ] (34), [Cu140 Se70 (PEt3 )36 ]
(36), and [Cu146 Se73 (PPh3 )30 ] (37) (without
organic groups). In 34 and 36, the two central
copper atoms are each half occupied to result
Fig. 3-44
in one atom in the sum, being statistically
distributed over two sites. Six further copper
atoms are assigned a 0.67 occupation,
resulting in four instead of six metal atoms in
the sum.
cluster framework in 33. With pseudo-C3h symmetry, the point group symmetry of
the molecule does not exceed C2 in the crystal because of slight distortions and the
positions of the organic groups bonded to phosphorus atoms.
The largest copper selenide clusters isolated to date possess 140 or 146 copper
atoms with 70 or 73 selenium ligands, respectively. The molecular structures of 34,
36 [23], and 37 are shown in Figure 3-44.
The different number of PR2 R 0 groups observed in 34 or 35 on the one hand and
36 on the other, affects the corresponding molecular structures in a manner similar to that for the ‘‘Cu70’’ species. Having a look at the corners again, 34, containing
thirty-four terminal PEt3 ligands, reveals an enlarged analog of 30. As in the
smaller compound, two corners are arranged irregularly, whereas the third one
shows the regular, centered cubic fragment. Increasing the number of PEt3 groups
to thirty-six in 36, all three corners show an ordered arrangement of copper, selenium and phosphorus atoms. This structural unit, which is only found for the
PEt3 -ligated ‘‘large’’ CuaSe clusters, underlines the structural relationship of 30,
34, and 36 and demonstrates another steric influence of certain types of PR2 R 0
molecules. All other phosphine ligands that surround ‘‘large’’ copper selenide
clusters feature larger cone angles, which could be the reason for a forced modification to a more staggered arrangement if four corner CuaPR2 R 0 units are present.
Consequently, in 35, all three corners show arrangements like that in the smaller
131
132
3 Syntheses and Characterizations
analog 31. Disordered copper atoms that increase the molecular pseudo-symmetry
are again present in the ‘‘Cu140’’ clusters. Additionally, there is a twofold disorder
of one central copper atom on two positions to result in an actual ‘‘Cu 0:5’’aSea
‘‘Cu 0:5’’ axis. Such metal atoms are again affected by a two-thirds occupancy factor
and lie in the plane spanned by the middle selenium layer. In addition, the six
next-but-one metal atoms viewed from the cluster center are disordered statistically
to amount to four atoms in the sum.
Three formal replacements [Se 2 Cu(PR2 R 0 2 )] of three corner selenium atoms removal of six PR2 R 0 molecules in 36, and full occupation of all split positions present in the ‘‘Cu140’’ clusters yield the composition and D3d symmetric molecular
structure of 37. Thus the structural transitions of the ‘‘Cu70’’ compounds to slightly
larger species and the transition occurring with the largest known copper selenide
clusters are related.
All of the ‘‘large’’ CuaSe clusters show an A-B-A packing of selenium hexagonal
layers, the metal centers occupying either trigonal or tetrahedral holes or being
situated outside the selenium layers in a likewise trigonal or tetrahedral coordination environment. Thus, the layer-type clusters clearly show a closer relationship to
solid-state structures than do the spherical ones. For the Cu2x Se phases, however,
only cubic and not hexagonal modifications are known. Interestingly, for the Cu2 S
(chalcosine) high temperature phase [38], a hexagonal lattice with mainly trigonally surrounded copper atoms in addition to tetrahedrally and linearly coordinated
ones does exist. In contrast, the cubic Cu2x Se crystal structures [39] can be derived from an anti-fluorite lattice [51] by positioning the selenium atoms on defined lattice places. Six to eight copper atoms per unit cell – according to the phase
composition – are localized in trigonal, tetrahedral or octahedral holes. In a-Cu2 Se,
the metal centers are placed in an ordered manner within the selenium lattice,
whereas in b-Cu2 Se they are statistically distributed. Even though the reported
structural properties of copper selenide binary phases are not uniform, there is
clear agreement regarding the dominating coordination number of four for the
copper atoms and the absence of any linearly coordinated metal centers. Because
of the hexagonal modification and the preferred coordination number of three for
the copper atoms, the layer-type CuaSe clusters approach the bulk material situation only in principle.
For the copper selenide clusters presented herein, one perceives that it is not
possible to assign a special range of cluster sizes to one given phosphine ligand,
being followed by another PR2 R 0 type to protect the next larger clusters and so on.
It can rather be recognized that even very different types of phosphines – perhaps
varying in size – are in the position to provoke the formation and crystallization of
similar, isomeric or even identical CuaSe cores if the ligand shells comply with
certain requirements of regular geometry and thus protective properties. In contrast, an examination of Scheme 3-3 rationalizes the formation of several clusters
that may vary significantly in size, in the presence of the same phosphine, where
different reaction conditions permit various suitable ligand shells to be realized. In
order to illustrate dependence of the size of the cluster core on the use of different
copper salts, stoichiometries, solvents, and reaction temperatures, the cluster for-
3.1 Semiconductor Nanoparticles
(15)
[Cu26Se13(PEt2Ph)14]
+ 1 PEt2Ph
–20 °C, Et 2O
CuOAc
+ 2 PEt2Ph
+ 20 °C, DME
CuCl
+
+
0.5 Se(SiMe3)2
0.5 Se(SiMe3)2
+ 1 PEt2Ph
(31)
[Cu70Se35(PEt2Ph)n] (n = 23, 24)
0 °C, Et2O
[Cu44Se22(PEt2Ph)18]
(23)
+ 1 PEt2Ph
0 °C, Diglyme
[Cu140Se70(PEt2Ph)34]
(35)
Scheme 3-5 Synthesis of various copper selenide clusters in
the presence of PEt2 Ph under different reaction conditions.
500
400
En [kJ·mol–1]
300
200
S
Se
100
1
10
5
n
Fig. 3-45 Course of the stabilization energies per monomer
unit En in ‘‘naked’’ clusters [Cu2n En ] (E ¼ S, Se; n ¼ 1–4, 6, 10).
The given values consider the total energies of the most stable
structural isomers computed on MP2 level.
mation reactions in the presence of PEt2 Ph as phosphine ligand are shown in
Scheme 3-5, with thorough description of the respective conditions.
As for the copper sulfide clusters, comprehensive theoretical investigations
within the MP2 approximation (program system TURBOMOLE) were carried out
for the [Cu2n Sen ] clusters [34]. Again, stabilization energies per monomeric unit
for a ‘‘naked’’ cluster growth from n ¼ 1 to n ¼ 6 and n ¼ 10 were calculated, as
were the phosphine binding energies for mono-, di-, tri-, tetra- and hexameric species. The results are given in Figures 18 and 19 respectively. The corresponding
graphs for the CuaS clusters (see Figures 5 and 6) are reproduced once more for
comparison.
133
134
3 Syntheses and Characterizations
500
S
Se
400
300
Eb [kJ·mol–1]
200
100
0
1
2
3
4
5
6
n
Fig. 3-46 Course of the CuaP binding energies
per CuaP bond Eb in phosphine-ligated clusters
[Cu2n En (PH3 )m ] (E ¼ S, Se; n ¼ 1–4, 6; m ¼ 2,
4, 6, 8). The given values consider the total
energies of the most stable structural isomers
computed on MP2 level with respect to the
total energy of PH3 calculated by the same
method.
The course of both energies is equivalent for the quantum chemically investigated sulfur- or selenium-bridged copper clusters. Thus the statement describing
the kinetic stability of the species and the very low CuaP binding energy is the
same as that for the CuaS compounds. Additionally, a comparison of the energetic
properties of sulfur- vs selenium-bridged copper clusters produces two results.
Firstly, with nearly identical courses of the stabilization energy (maximum differences: G5 kJmol1 , <2%) up to n ¼ 4, one notices a stabilization gain for the
CuaS species from n ¼ 6 onwards. This correlates with a larger absolute value for
the formation enthalpy of solid Cu2 S (DHf : 83.4 kJmol1 ) [49] when compared
to solid Cu2 Se (DHf : 59.3 kJmol1 ) [50]. The stabilization energies per monomeric unit for the calculated, rather small, clusters however do not exceed differences of 5%. Thus it is comprehensible why laser ablation spectra of copper chalcogenides do not vary much for different chalcogen types. Secondly, the phosphine
molecules are generally bonded by 4–7 kJmol1 (4–9%) higher values of the CuaP
binding energy for the sulfur-bridged compounds. This correlates with the discrepancy between the experiments reported by Dance et al. [32] that show similar
behavior for different chalcogens – in agreement with the calculations of ‘‘naked’’
species – and the experimental observations of significant differences in the syntheses and structures of phosphine-clad sulfur or selenium-bridged copper clusters, respectively. The larger CuaP binding energies computed for CuaS clusters,
when compared to the CuaP binding energies of their CuaSe analogs help to explain the different synthetic product spectra. Assuming that the activation energy
Ea for CuaP bond cleavage increases by a value comparable to the CuaP binding
energy itself when going from the CuaSe to the CuaS systems, one calculates a
decrease in the respective reaction rate constant k by means of the Arrhenius
equation (Eq. 1) [52].
3.1 Semiconductor Nanoparticles
k ¼ A eEa =ðRTÞ
For a typical cluster formation temperature T, e.g. 250 K, an increase in the activation energy Ea of about 5 kJmol1 causes a reduction in the rate constant k by one
order of magnitude. Thus, the cluster growth that is initiated by the descent of a
primary terminal phosphine ligand shell will occur less intensely for sulfur-bridged
clusters. Consequently, unlike the selenium-bridged copper clusters, the CuaS
species show a much lower tendency to exceed the aggregation beyond ‘‘Cu12’’.
The use of alkylated or arylated derivatives RSeSiMe3 (R ¼ organic group) provokes syntheses according to another mechanism that leads to the formation of
selenido-/selenolato-bridged copper clusters. The given synthesis route (Schemes
3-1 and 3-4) in no case produces selenolato-free species. This differs from similar
reactions with tellurolato reactants RTeSiMe3 (see below). The experimentally isolated and characterized compounds that are listed in Scheme 3-4 display significantly different structures with regard to the purely E 2 -bridged molecules (E ¼ S,
Se). This applies even for reactions in the presence of tertiary phosphines PPh3 or
PiPr3 that were likewise used for the syntheses of copper selenide clusters. The
expanded ligand properties of SeR , which can act as both a bridging and a terminal ligand, leads to the crystallization of low nuclearity complexes with two,
three, six or nine metal centers respectively (38–42) [20, 22, 30] if SeaC bond
cleavage is suppressed. Larger cluster compounds can be isolated whenever Se 2
particles are produced during the reaction to make inner selenium bridges available. However, these molecules containing selenide and selenolato ligands show
structural patterns that are also common for CuaSe clusters as far as the inner
core is concerned. In Figure 3-47, the compounds illustrated (38–42) contain only
SeR and no Se 2 ligands.
Binuclear complexes 38 [30] and 39 [20] formally show analogous compositions,
but structurally differ according to the restrictions of the different types of ligating
components present. These restrictions have structural relevancy for all Se 2 /
SeR -bridged copper clusters as follows. Firstly, bidentate phosphines dppR
(R ¼ organic spacer) bridge two copper centers at the periphery of the polynuclear
complex or cluster molecule and therefore simultanously act as protecting groups.
Secondly, selenolato ligands SeR usually behave like dppR molecules although
they are able to bridge two or three metal centers. Additionally, they can be terminally bonded, as found in 39 when the CuaCu distance (9.01 Å) is too large for a
m2 -SeR bridge. Thirdly, tertiary phosphines PR2 R 0 act exclusively as terminal ligands, and thus the copper atoms in 38 are held by SePh fragments at a distance
of 2.846(1) Å. The molecular structure of 40 [22] also follows these patterns, again
with the final purpose of optimal protection of the heavy atom framework by organic groups. Compounds 41 [30] and 42 [30] are the only examples that could be
isolated by employing the bidentate amine ligand bipyridyl (41) or diphenylphosphino acid Ph2 PO2 H (42). Cluster 41 seems to represent a ‘‘snapshot’’ molecule of
cluster growth, since two [Cu(bipy)]þ fragments are unsymmetrically attached to a
quite regular cluster center. The latter is formed from a distorted tetrahedron Cu 4
(CuaCu: 2.650(4)–2.824(4) Å) whose six edges are m2 -bridged by SePh ligands.
135
136
3 Syntheses and Characterizations
39
38
40
41
42
Molecular structures of [Cu2 (SePh)2 (PPh3 )3 ] (38),
[Cu2 (SePh)2 (dppbd)3 ] (39), [Cu3 (SeMes)3 (dppm)] (40),
[Cu 6 (SePh)6 (bipy)2 ] (41), and [Cu9 (SePh)6 (O2 PPh2 )3 ] (42)
(without hydrogen atoms).
Fig. 3-47
Four of the phenyl selenolato groups, however, act as m3 bridges, since each pair
binds to the additional Cu(bipy)þ units. Much higher molecular and crystallographic symmetry is found for 42. Two six-membered Cu3 (SePh)3 rings are conjunct by three ‘‘naked’’ copper centers leading to a m3 bridging function of the
SePh groups. Additionally, the rings are held by three Ph2 PO2 ligands that bind
to the six copper atoms. Twelve phenyl groups envelop the uncharged cluster and
therefore enable the termination of atomic aggregation and crystallization in THF.
As for the copper chalcogenide clusters 1–37, the nature of the phosphine ligands plays a key role in the size and structures of the observed product molecules
– even though they are no longer the only protective groups at the cluster peripheries. Success in the synthesis of copper selenide/selenolate clusters was achieved
by the use of bidentate phosphines. The variation of organic spacers between the
two ligating [Ph2 P] fragments replaces the former variation of organic substituents
R of PR2 R 0 with the selenolato-free clusters. Again, one given dppR ligand does
not necessarily restrict the synthesis to the formation of one single cluster product.
Outer reaction conditions rather select between several potential compounds that
can all be conveniently protected by the phosphine employed. Thus, with ethane,
propane, butane, acetylene or benzene spacers, one obtains clusters with 16, 25,
36, 38, or 58 copper centers (43 [20], 44 [22], 46–50 [22, 30], 52 [22]). The only
tertiary phosphine PR3 that was successfully used for the formation of higher nuclearity Se 2 /SePh -bridged copper clusters is PiPr3. The synthesis of molecules
featuring thirty-two, fifty, or even seventy-three metal centers (45, 51, 53) [23] was
observed by variation of the reaction conditions with this phosphine. Figure 3-48
3.1 Semiconductor Nanoparticles
43
44
Molecular structures of [Cu16 Se4 (SePh)8 (dppbe)4 ]
(43) and [Cu(dppp)2 ][Cu25 Se4 (SePh)18 (dppp)2 ] (44) (without
terminal phenyl groups).
Fig. 3-48
shows the molecular structures of 43 and the anion in 44, the only ionic copper
selenide cluster compound isolated to date via this synthetic method.
Both clusters 43 and 44 are oblong in shape. The selenium substructure of 43
can, however, be viewed as a beginning of hexagonal packing of selenium atoms.
The middle ‘‘layer’’ consists of six atoms arranged in zigzag fashion, the two
enclosing ‘‘layers’’ each containing three selenium centers occupying the quasi A
position with respect to the quasi B position of the atoms in between. Four of the
chalcogen (Se 2 ) bridges act as m4 or m6 bridging ligands. The SePh ligands are
bonded to two or three copper neighbors each and at the same time contribute to
the protection of the CuaSe core by means of their organic substituents. Six copper
atoms are surrounded linearly by two selenium ligands: two show an almost
trigonal planar [Se3 Cu] coordination, whereas the remaining eight metal centers
occupy the centers of slightly distorted [Se3 P] tetrahedra. In contrast to 43, the
chalcogen framework of 44 rather represents an ellipsoid-type deltahedron of four
Se 2 ligands and 18 selenium atoms of the SePh groups. Again, the bridging
grade of the chalcogen centers is higher for the ‘‘naked’’ ligands (m5 or m7 ) than for
the SePh fragments (m2 –m4 with regard to the CuaSe bonds). There are eighteen
copper atoms positioned at the cluster periphery, and these belong to almost planar
[Se3 Cu] or distorted tetrahedral [Se3 CuP] units. The observation that no linear
[Se 2 Cu] arrangements are present in 44 underlines the trend to higher coordina-
137
138
3 Syntheses and Characterizations
45
46
Molecular structures of [Cu32 Se7 (SenBu)18 (PiPr3 )6 ]
(45) and [Cu36 Se5 (SePh)26 (dppa)4 ] (46) (without n-butyl, ipropyl or phenyl groups). Three copper atoms (marked with a
cross) in the cluster center of 46 are distributed over six sites
with an occupation of 0.5.
Fig. 3-49
tion numbers at the metal center with increasing molecular size as reported for the
copper selenide clusters without selenolato ligands. The cluster corpus is large
enough to accommodate some inner copper atoms that are either two-, three-, or –
rarely – five-coordinated by selenium neighbors. Comparing the compositions of
43 and the anion in 44, one perceives an increasing preference for the ligation by
SePh ligands rather than the ligation by phosphine donors. This may be explained by the ambivalent character of the SePh groups, which firstly form fairly
stable CuaSe cluster bonds and secondly carry protective organic substituens.
Thus, the ratio of ‘‘terminally ligating selenium atoms’’ to phosphorus atoms increases from 1:1 in 43 to 4.5:1 in 44.
The molecular structures of the cluster compounds 45 and 46 are given in
Figure 3-49. Figure 3-50 shows the packing of selenium atoms in 45 and the selenium substructure in 46.
As for the comparison of 43 and 44, one can discuss a smaller cluster whose
structure is based on a regular packing of selenium atoms in layers (45) and a
larger, rather spherically shaped compound, that features a polyhedral selenium
framework (46). The selenium layers in the SenBu -bridged species 45 are packed
3.1 Semiconductor Nanoparticles
Selenium substructures in 45 (top) and 46 (bottom).
The drawn SeaSe contacts do not represent bonding interactions, but only serve to visualize the A-B-C packing of selenium
atoms in three layers in 45 or explain the geometry of the
polyhedron in 46.
Fig. 3-50
in an A-B-C fashion similar to the Cu2 Se solid state structure, which is so far unparalleled for all known copper chalcogenide clusters. This similarity between the
atomic arrangement of a cluster molecule and the non-molecular Cu2 Se suggests
that a fragment of a bulk structure may indeed be stabilized to nano-size by coating the surface with suitable groups. This tendency is also observed in three silver
selenide/selenolate clusters, [Ag112 Se32 (SenBu)48 (PtBu3 )12 ], [Ag114 Se34 (SenBu)46 (PtBu3 )14 ], and [Ag172 Se40 (SenBu)92 (dppp)4 ] (see section 3.1.3.3) [53]. Cluster 45 is
the only copper selenide/selenolate compound that contains SenBu groups like
these silver clusters. It is not yet proven, but it is at least possible, that the packing
of the selenium atoms may be influenced by the nature of the organic group attached to the selenolato ligand. The chalcogen layers in 45 are centered by a large
planar rhombus of nine selenium atoms. Five of these are ‘‘naked’’ Se 2 ligands in
the cluster network, while the four vertex atoms are SenBu groups. Above and
below this middle layer, one finds two truncated rhombuses consisting of only
eight selenium atoms, the ‘‘central’’ ones not being bonded to organic substituents.
However, the size dependency of structural properties seems to be different from
that of the selenolato-free copper clusters: layer-type structures are observed even
for smaller molecules like 43 or 45 in addition to larger, spherical ones. A reversion
to lower average coordination numbers in larger clusters may also appear, as observed in 45. Of the 32 copper atoms in 45, twelve have a trigonal planar coordination geometry and two show an ideal tetrahedral environment. The remaining
139
140
3 Syntheses and Characterizations
47
49
Molecular structures of [Cu36 Se13 (SePh)12 (dppe)6 ]
(47) and [Cu38 Se13 (SePh)12 (dppp)6 ] (49) (without phenyl
groups).
Fig. 3-51
eighteen metal atoms are shifted away from the centers of tetrahedral holes.
Hence, each of them displays only three normal CuaSe bonds, with the fourth
CuaSe distance being relatively long (3.18–3.35 Å). The tendency to occupy tetrahedral holes – even without perfect realization – again correlates with the structural properties of bulk Cu2 Se. In 46, five inner Se 2 ligands are enclosed by a
shell of twenty-six SePh fragments. The bridging grades of the Se 2 ligands
amount to m7 or m8 whereas the selenium atoms of the SePh groups act only as
m2 –m4 bridges between next copper atoms. The copper centers feature coordination
numbers of two, three, or four, within [Se 2 Cu], [Se3 Cu], [Se4 Cu], or [Se3 CuP]
units, respectively. Depending on their position inside the cluster or at the surface,
the selenium neighbors of the metal atoms are either exclusively Se 2 or a combination of Se 2 and SePh ligands. As for the 43 and 44, the number of terminal
selenolato groups is a multiple of the number of ligating phosphorus atoms (SeR:P
ratios are 3:1 in 45, 3.25:1 in 46).
Four closely related CuaSeaSePh clusters are formed in the presence of either
dppe, dppp, or dppb (47–50). The syntheses of isomeric, mixed valence compounds 47 and 48 containing 36 copper centers were carried out in THF; the isomeric ‘‘Cu38’’ clusters 49 and 50 that formally display exclusively Cuþ crystallize
from toluene. As well as being performed in different solvents, the reactions producing the dppp-ligated compounds 48 or 50 differed in the CuCl:dppp ratios
(3.1:1 for 48, 5:1 for 50). In Figure 3-51, the molecular structures of 47 and 49 are
given as examples.
Cluster 47 crystallizes in the trigonal space group R3, whereas 48–50 crystallize
in triclinic space (P1). Despite the different crystallographic symmetry, all four
compounds display very similar structural properties. A closer look at the atomic
arrangements helps to put the preference of this prototypic cluster framework
3.1 Semiconductor Nanoparticles
Fig. 3-52 Central [Se13 Cu12 ] unit in 47. The drawn lines only
represent polyhedral edges and do not specify bonding
interactions.
down to highly symmetric substructures. A central selenium atom positioned on
the inversion centers of the clusters is surrounded by a distorted icosahedron of
twelve copper atoms. The latter is itself enclosed by another distorted Se12 icosahedron. An alternate packing of Cu12 and Se12 icosahedra has already been discussed for the CuaSe clusters 14 and 15 (Figures 3-34 and 3-35). In these compounds, one polyhedron is positioned with its edges perpendicular to the edges of
the polyhedron below. In contrast to this pattern, the Se12 and Cu12 icosahedra
in 47–50 are arranged such that one of the idealized fivefold axes of one polyhedron parallels one of the idealized threefold axes of the other and vice versa. The
central [Se13 Cu12 ] unit of 47 is shown as an example in Figure 3-52.
In this way, six edges and six faces of the Cu12 icosahedron are bridged by the
selenide ligands that form the Se12 polyhedron. The cluster core expands out to the
periphery by coordination of the Se12 substructure to another 24 copper atoms,
twelve of which are bonded to the six dppR ligands. These copper atoms are
bridged by twelve m3 -SePh groups that are positioned at the cluster surface with
the phosphine molecules. Except for the m12 bridging central selenium atom, the
selenide ligands act as m6 (each six in 47 and 48, all in 49 and 50) or m5 bridges
(each six in 47 and 48) between copper centers. The different numbers of copper
neighbors per selenium atom in 47 and 48 on the one hand and 49 and 50 on the
other hand is caused by two copper atoms being formally ‘‘missing’’ in the molecular structures of 47 or 48, resulting in the mixed-valence situation. The metal
centers are disorted trigonal planar or tetrahedrally surrounded. All CuaSe distances in the mixed-valence clusters lie within the range of the CuaSe bond
lengths in all four clusters (2.301(4)–2.888(3) Å) and thus give no evidence of any
localization of the Cu 2þ centers. However, the range of the CuaCu contacts in 47
or 48 (2.511(3)–2.943(4) Å) is somewhat narrower than that found in 49 or 50
141
142
3 Syntheses and Characterizations
(45)
[Cu32Se7(SenBu)18(PiPr3)6]
m = 0.53
0°C → +20°C
Et2O
+ 1.2 nBuSeSiMe3
CuCl
+
m PiPr3
m=1
+ 1.0 tBuSeSiMe3
m=2
+ 0.5 PhSeSiMe 3
+ 20 °C, Et 2O
– 70 °C, Et2O
+ Se(SiMe 3)2
→ + 20 °C
[Cu73Se35(SePh)3(PiPr3)21]
(53)
[Cu50Se20(SetBu)10(PiPr3)10]
(51)
Synthesis of various selenido/selenolato-bridged
copper clusters in the presence of PiPr3 under different reaction
conditions.
Scheme 3-6
51
Fig. 3-53
Molecular structure of [Cu50 Se20 (SetBu)10 (PiPr3 )10 ] (51) (without organic groups).
(2.305(4)–3.040(3) Å). The spherical, somewhat oblate molecules show maximum
diameters of 23.9–25.7 Å.
The use of different selenolato sources and different optimized reaction conditions leads to three various copper selenide/selenolate clusters with terminal
PiPr3 ligands, 45 (Figures 22 and 23), 51, and 53. The latter is synthesized in an
uncommon manner by the use of both PhSeSiMe3 and Se(SiMe3 )2 in the same
reaction and is discussed below. Scheme 3-6 summarizes the syntheses of the
PiPr3-clad clusters, specifying the reaction conditions in detail, and Figure 3-53
shows the molecular structure of compound 51.
Examining Figure 3-53, one surmises that a ‘‘normal’’ cluster growth occurred
in the presence of a tertiary phosphine – indeed under SeaC bond cleavage
3.1 Semiconductor Nanoparticles
52
Fig. 3-54
Molecular structure of [Cu58 Se16 (SePh)24 (dppa)6 ] (52) (without phenyl groups).
conditions – to result in the formation of a copper selenide cluster which is represented by the [Cu40 Se20 ] central unit of 51. Only the sidewise attachment of two
symmetry equivalent [Cu5 (SetBu)5 ] fragments reflects the presence of the selenolato reactant. This is in contrast to the related PiPr3-ligated structure of 45, which
shows only seven Se 2 ligands with eighteen SetBu groups. A definite correlation
between the significantly different course of the cluster formation and the respective reaction conditions cannot be evaluated based on the low statistic, but it is very
likely that the Cu:P ratio of the reactants and the temperature play an important
role. The copper atoms in 51 bind to Se 2 , SePh , and PiPr3 ligands in different
ratios to achieve coordination numbers of two (with a bend coordination geometry
indicating a weak third interaction), three, and four. The selenium atoms that carry
t-butyl groups act as m3 bridges between copper atoms, whereas the Se 2 ligands
bridge between five, six or seven copper neighbors. The molecules are regularly
protected by PiPr3 ligands and the organic substituents of the SetBu groups.
The molecular structure of another dppR-ligated copper cluster (R ¼ a ¼
acetylene) with both selenido and selenolato ligands, 52, again displays high symmetry, and here the molecular symmetry is underlined by very high crystallographic symmetry. The compound crystallizes in the cubic space group F43c and
is thus the second copper chalcogenide cluster known besides 14 (space group
Fm3) with a cubic crystal lattice. The molecular structure of 52 is illustrated in
Figure 3-54.
As observed for 47–50, the cluster core is oblate in shape. However, in contrast
to 47–50, it is not based on icosahedral fragments. Instead, one perceives a central
tetrahedron Cu 4 sharing its corners with four further Cu 4 units that themselves
share the trans face with each octahedral Cu 6 . The remaining 30 copper atoms are
143
144
3 Syntheses and Characterizations
Fig. 3-55 Copper substructure in 52. The connected tetrahedra
Cu 4 and octahedra Cu 6 are highlighted by gray shading. CuaCu
distances are drawn up to 301.2(4) Å.
situated between these polyhedra or are connected at the outside of the polyhedral
arrangement. The copper substructure of 52 is given in Figure 3-55.
Even though the Cu58 substructure has to be assigned the formal charge 56þ
(assuming Se 2 and SePh ligands as well as electroneutrality in the sum), which
leads to an extraordinary mixed valence Cuþ /Cu 0 compound, one does not find
any hints of exceptional CuaCu binding interactions. The CuaCu distances lie
within the usual range of 2.540(3) and 3.012(4) Å. Thus, the cluster core is habitually held together by bridging selenido or selenolato ligands. The former act as m5
or m6 bridges, and the latter are m3 bridges between next neighboring copper atoms.
Despite the large cluster size, near-linear as well as trigonal planar or tetrahedrally
coordinated metal centers are observed. Therefore, it is again proved that the sizedependence of the structural principles discussed for chalcogenolato-free copper
sulfide or copper selenide clusters are widely invalid in the CuaSeaSeR cluster
systems. Six dppa ligands and 24 phenyl groups of the SePh fragments ligate the
cluster molecule. Hence 52 continues the series of compounds that show a significant preference for protection by SeR groups rather than phosphine ligands. The
clusters 43 and 47–51 represent exceptions to this series, as their SeR:P ratios are
all 1:1.
The largest copper cluster with both Se 2 and SeR ligands in the framework is
53. Its molecular structure is shown in Figure 3-56.
The cluster displays a [Cu70 Se35 ] core that is identical to that observed in 28 or
29 (Figures 15 and 16). Only three additional CuaSePh fragments positioned at the
corners of the triangular molecule distinguish 53 from the slightly smaller clusters. The fact that one observes a typical copper selenide cluster structure with only
a few selenolato groups present is surprising, since PhSeSiMe3 was the first sele-
3.1 Semiconductor Nanoparticles
53
Fig. 3-56
Molecular structure of [Cu73 Se35 (SePh)3 (PiPr3 )21 ] (53) (without organic groups).
nium source added during the synthesis. This indicates again the distinct preference for the ‘‘Cu70’’ molecular structure that is formed whenever the reaction conditions (Scheme 3-6) and the solubility situation enables the cluster growth to proceed until crystallization of this large cluster compound can occur.
Tellurium-Bridged Copper Clusters
In this section, we focus on the description of the syntheses and structures of
ligand-stabilized copper telluride cluster molecules. Solid-state copper tellurides as
well as ternary alkali metal copper telluride compounds, which consist of ndimensionally linked copper telluride clusters, are beyond the scope of this review.
Their syntheses have accelerated lately, although only a limited number exist up to
now, e.g., KCuTe [54], NaCuTe [55], KCu3 Te 2 [56], NaCu3 Te 2 [57], K2 Cu5 Te5 [58],
and K 4 Cu8 Te11 [59], all of which have been prepared by means of melt reactions.
The recently synthesized compound K3 Cu11 Te16 [60] is the first example of a ternary alkali metal copper telluride crystallized from supercritical 1,2-diaminoethane.
According to the general reaction pathway given in Scheme 3-7 and 3-8, the
syntheses of ligand-stabilized copper telluride clusters have mainly been achieved
in one of three different ways [11, 16, 19, 20, 24, 25, 31, 61].
Reaction of copper(I) acetate with Te(SiMe3 )2 at low temperatures (50 to
20 C) yields smaller cluster compounds with stoichiometric composition. At
higher reaction temperatures, up to 25 C, larger clusters can be isolated that
are mostly of the mixed-valence type. Similar mixed-valence compounds have
also been isolated from reactions of copper(I) chloride with Te(SiMe3 )2 at lower
145
146
3 Syntheses and Characterizations
CuOAc + 0.5 Te(SiMe3)2
CuCl + 0.5 Te(SiMe3)2
CuOAc + nBuTe(SiMe3)2
CuCl + nBuTe(SiMe3)2
CuCl + tBuTe(SiMe3)2
CuCl + PhTe(SiMe3)2
+ PR3 or PR2R´
– SiMe3OAc
+ PR3 or PR2R´
– SiMe3Cl
+ PR3 or PR2R´
– SiMe3OAc
+ PR3 or PR2R´
– SiMe3Cl
+ PR3 or PR2R´
– SiMe3Cl
+ PR3 or PR2R´
– SiMe3Cl
[Cu4Te(Te2)2(PCyPh2)4]
[Cu6Te3(PCyPh2)5]
[Cu8Te4(PPh3)7]
[Cu12Te6(PPh3)8]
[Cu16Te8(PPhnPr2)10]
[Cu16Te9(PPh3)8]
[Cu23Te13(PPh3)10]
[Cu24Te12(PiPr3)12]
[Cu26Te12(PEt2Ph)12]
[Cu44Te23(PPh3)15]
[Cu44Te23(PPhnPr2)15]
[Cu44Te23(PCyPh2)15]
[Cu58Te32(PCy3)16]
(59)
(54)
(55)
(56)
(57)
(73)
(76)
(58)
(79)
(82)
(83)
(84)
(85)
[Cu4(Te2)2(PiPr)4]
[Cu16Te9(PEt3)8]
[Cu16Te9(PEt2Ph)8]
[Cu23Te13(PiPr3)10]
[Cu26Te10(Te2)3(PtBu3)10]
[Cu28Te13(Te2)2(PEt2Ph)12]
[Cu29Te16(PiPr3)12]
(60)
(74)
(75)
(77)
(61)
(62)
(81)
[Cu11Te(TenBu)9(PPh3)5]
(65)
[Cu18Te6(TenBu)6(PnPr3)8]
[Cu23Te13(PEt3)12]
[Cu58Te32(PPh3)16]
(69)
(78)
(86)
[Cu18Te6(TetBu)6(PEtPh2)7]
[Cu19Te6(TetBu)7(PEt3)8]
[Cu27Te15(PMeiPr2)12]
[Cu58Te32(PnButBu2)14]
(68)
(70)
(80)
(87)
[Cu6(TePh)6(PEtPh2)5]
[Cu12Te3(TePh)6(PPh3)6]
[Cu12Te3(TePh)6(PEt3)6]
[Cu5(TePh)6(PEt)3][PEt3Ph]
(64)
(66)
(67)
(63)
Survey of the reaction pathways for the synthesis
of ligand-stabilized tellurido or tellurido/tellurolato-bridged
copper clusters.
Scheme 3-7
[Cu12Te3(TePh)6(PEt3)6] (67)
[Cu5(TePh)6(PEt)3][PEt3Ph] (63)
[Cu6(TePh)6(PEtPh2)5] (64)
THF / RT
− 1,5 TePh2
hn / C6H6
− TePh2
[Cu29(TePh)12Te9(PEt3)8] [PEt3Ph]
[Cu50Te17(TePh)20(PEtPh2)8][PEtPh3]4 (72)
Reaction pathways for the formation of larger
telluride/tellurolato-bridged copper clusters starting from
cluster precursors.
Scheme 3-8
(71)
3.1 Semiconductor Nanoparticles
temperatures (down to 70 C). By treating copper(I) chloride or copper(I) acetate with tellurolate compounds RTeSiMe3 (R ¼ tBu, nBu, Ph), one obtains clusters which contain both RTe and Te 2 ligands as a result of the facile cleavage of
the TeaC bond. In some cases, we isolated clusters featuring TeaTe bridges. Examples of syntheses of larger copper telluride clusters starting from cluster precursors are light-induced formation of [Cu50 Te17 (TePh)20 (PEtPh2 )8 ][PEtPh3 ]4 (72)
from [Cu 6 (TePh)6 (PEtPh2 )5 ] (64) in benzene [19] and the co-condensation of
[Cu12 Te3 (TePh)6 (PEt3 )6 ] (67) with [Cu5 (TePh)6 (PEt)3 ][PEt3 Ph] (63) to form [Cu29 (TePh)12 Te9 (PEt3 )8 ] [PEt3 Ph] (71) as outlined in Scheme 3-8 [61].
Table 3-1 summarizes all the CuaTe cluster compounds that have been synthesized and structurally characterized by single-crystal X-ray diffraction analysis to
date. The species can be subdivided into four main groups, three of which consist
of stoichiometric compositions [Cu2n Tenx (TeR)2x (PR2 R 0 )m ] that differ in the nature of their tellurium ligands, while the fourth represents the group of mixedvalence clusters.
The structures of the tellurium-bridged copper clusters (Figures 3-57–3-61, see
below) differ from those of the copper sulfide or selenide clusters, with the exception
of some smaller ones. This can be understood against the background of the larger
atomic and ionic radii of the tellurium atoms compared to the lighter chalcogen
atoms, allowing higher coordination numbers and therefore new binding geometries which lead to different structures. In addition, one finds a greater tendency
to form mixed-valence clusters. This is in line with the existence of many stable
binary compounds with compositions Cu2x Te [40] and with the situation found in
the ternary alkali copper telluride compound K2 Cu5 Te5 [58].
In the case of copper telluride clusters, we have not yet reached the turning point
in size where the whole core structures display bulk structure characteristics, as
seen for the copper selenide species. Nevertheless, the tellurium frameworks in the
largest cluster compounds, 85–87 display hexagonal structure properties. Powder
diffraction patterns of Cu2 Te could be also indexed on the basis of a hexagonal cell
of tellurium atoms. However, the structural details for bulk Cu2 Te have not yet
been completely established [40b].
In all the copper tellurium cluster molecules, the tellurium atoms form polyhedra with triangular faces. In the larger ones, some of the tellurium atoms are
located inside the polyhedra. For smaller cluster molecules, the tellurium polyhedra can usually be derived from classical polyhedra (Figure 3-57). Those of the
larger ones often show unusual cage structures. The distances between the tellurium atoms are usually non-bonding, except the TeaTe units in 59–62.
The copper atoms reside in the holes and on the surface of the polyhedra and are
coordinated by either two, three, or four tellurium atoms, with some being additionally coordinated by one further phosphine ligand.
The CuaTe distance ranges that are regarded as bonding contacts range from
2.403 to 3.337 Å for all the clusters under discussion. Even though coordination
numbers (CN) for Te 2 ligands range from 4 to 12 and those for TeR are usually
smaller (CN ¼ 2, 3, 4), there is no significant difference between the average
CuaTe distances for Te 2 and TeR ligands.
147
148
Classification of copper telluride into stoichiometric or mixed valence according to the types of tellurium ligands.
stoichiometric
only with Te 2C ligands
with additional TexTe bridges
with additional TeRC ligands
mixed-valence
[Cu 6 Te3 (PCyPh2 )5 ]
(54) ½31 [Cu 4 Te(Te 2 )(PCyPh2 )4 ]
(59) ½31 [Cu5 (TePh)6 (PEt)3 ][PEt3 Ph]
(63) ½61 [Cu16 Te9 (PPh3 )8 ]
[Cu16 Te9 (PEt3 )8 ]
[Cu16 Te9 (PEt2 Ph)8 ]
(73) ½24
(74) ½11
(75) ½11
[Cu8 Te4 (PPh3 )7 ]
(55) ½24 [Cu 4 (Te 2 )2 (PiPr3 )4 ]
(60) ½11 [Cu 6 (TePh)6 (PEtPh2 )5 ]
(64) ½19 [Cu23 Te13 (PPh3 )10 ]
[Cu23 Te13 (PiPr3 )10 ]
[Cu23 Te13 (PEt3 )12 ]
(76) ½24
(77) ½11
(78) ½16
[Cu12 Te6 (PPh3 )8 ]
(56) ½24 [Cu26 Te10 (Te 2 )3 (PtBu3 )10 ]
(61) ½11 [Cu11 Te(TenBu)9 (PPh3 )5 ]
(65) ½16 [Cu26 Te12 (PEt2 Ph)12 ]
(79) ½24
[Cu16 Te8 (PPhnPr2 )10 ] (57) ½24 [Cu28 Te13 (Te 2 )2 (PEt2 Ph)12 ] (62) ½11 [Cu12 Te3 (TePh)6 (PPh3 )6 ]
[Cu12 Te3 (TePh)6 (PEt3 )6 ]
(66) ½19 [Cu27 Te15 (PMeiPr2 )12 ]
(67) ½61
(80) ½25
[Cu18 Te6 (TetBu)6 (PEtPh2 )7 ]
[Cu18 Te6 (TenBu)6 (PnPr3 )8 ]
(68) ½25 [Cu29 Te16 (PiPr3 )12 ]
(69) ½16
(81) ½11
[Cu19 Te6 (TetBu)7 (PEt3 )8 ]
(70) ½25 [Cu44 Te23 (PPh3 )15 ]
[Cu44 Te23 (PPhnPr2 )15 ]
[Cu44 Te23 (PCyPh2 )15 ]
(82) ½24
(83) ½24
(84) ½24
[Cu29 (TePh)12 Te9 (PEt3 )8 ][PEt3 Ph]
(71) ½61 [Cu58 Te32 (PCy3 )16 ]
(85) ½31
[Cu58 Te32 (PPh3 )16 ]
(86) ½16
[Cu58 Te32 (PnButBu2 )14 ] (87) ½25
[Cu24 Te12 (PiPr3 )12 ]
(58) ½24
[Cu50 Te17 (TePh)20 (PEtPh2 )8 ][PEtPh3 ]4 (72) ½19
3 Syntheses and Characterizations
Tab. 3-1
Te6 56
Te8 57
Te12 58
Te12(Te1) 76, 77, 78
Te12 79
Te21(Te2) 82, 83, 84
Te15(Te1) 61
Te28(Te4) 85, 86, 87
149
Tellurium polyhedra in copper telluride clusters.
Connections to the central tellurium atoms are omitted for
clarity. Direct TeaTe bonds in 61 and 62 are drawn thicker.
Fig. 3-57
Te15(Te1) 81
3.1 Semiconductor Nanoparticles
Te16(Te1) 62
Te14(Te1) 80
Te9 73, 74, 75
150
3 Syntheses and Characterizations
Those tellurium atoms that are located in the center of the cluster molecules
tend to bond to a large number of copper atoms. In some compounds, the copper
atoms themselves form polyhedra, which can be described as Frank-Kasper polyhedra [62]. In these cases, the clusters consist of interpenetrating Frank-Kasper
polyhedra, which are well known from a number of intermetallic phases (e.g.,
Laves Phases).
However, the observed CuaCu distances give no indication of strong metal-metal
(i.e. d 10 –d 10 ) interactions. Theoretical investigations point to the same conclusions
[34b, 63]. The shortest CuaCu contacts were observed in the molecular structures
of 62 (2.383 Å), and 78 (2.382 Å), but these involve copper atoms which exhibit
elongated thermal ellipsoids. Aside from these, the shortest CuaCu distances are
interestingly observed for the largest clusters, namely 85 (2.426 Å), 86 (2.448 Å),
and 87 (2.433 Å).
In the following section, all cluster structures will be briefly described and discussed according to the division into four groups indicated in Table 3-1.
The molecular structures of the stoichiometric copper telluride clusters 54–58
are shown in Figure 3-58.
54
56
55
57
58
Molecular structures of stoichiometric copper
telluride clusters [Cu 6 Te3 (PCyPh2 )5 ] (54), [Cu8 Te4 (PPh3 )7 ] (55),
[Cu12 Te6 (PPh3 )8 ] (56), [Cu16 Te8 (PPhnPr2 )10 ] (57), and
[Cu24 Te12 (PiPr3 )12 ] (58). C and H atoms are omitted for clarity.
Fig. 3-58
3.1 Semiconductor Nanoparticles
In 54, three tellurium atoms form a triangle, in the center of which a copper
atom is coordinated in a slightly distorted triangular fashion. Four [CuPPh2 Cy]
units are doubly bridging two of the edges. The third edge is bridged by a copper
atom which is bonded to two phosphine ligands.
The tellurium atoms in 55 form a tetrahedron, the six edges of which are
bridged by [CuPPh3 ] groups. An additional m3 -[CuPPh3 ] fragment is bonded to a
Te3 face, and one naked copper atom is located in the center of the Te4 tetrahedron.
Cluster 56 features a Te6 octahedron. Six [CuPPh3 ] groups act as m2 -bridges
above six of the octahedral edges. Two further [CuPPh3 ] groups are coordinated by
three tellurium atoms, giving a tetrahedral environment around the copper atoms.
The other four copper atoms are each coordinated by three tellurium atoms. Thus,
56 represents a novel isomer of the [Cu12 E6 (PR3 )8 ] clusters (see above) which was
not observed for E ¼ S, Se.
An octahedron of tellurium atoms is also found in 57. However, two further tellurium atoms additionally cap two opposite triangular faces. Eight [CuPPhnPr2 ]
groups are bonded along edges of this polyhedron. Of the other eight copper centers, two are tetrahedrally coordinated by one phosphine and a Te3 face. The six
remaining metal atoms are bonded in a distorted trigonal planar mode by three
tellurium atoms and are located inside the cluster cavity.
The Te12 substructure of 58 cannot be described as a classical deltahedron, but
rather shows an arrangement of two highly distorted face-sharing tetragonal antiprisms. Four [CuPiPr3 ] groups bridge one TeaTe non-bonding edge each, and
eight [CuPiPr3 ] units act as m3 -bridges above Te3 faces. The other twelve copper
atoms are each surrounded by three tellurium atoms and are slightly shifted away
from the Te3 planes toward the center of the molecule.
Some reactions of copper(I) chloride or copper(I) acetate lead to the formation of
copper telluride clusters that contain TeaTe bridges (59–62, Figure 3-59).
In 59, three tellurium atoms form an approximate isosceles triangle with a
TeaTe bond along the short edge (TeaTe: 2.791 Å). The two longer edges of the triangle are bridged by one [CuPPh2 Cy] unit each. The two other copper atoms are
located above and below the triangular face and are additionally coordinated by
phosphine ligands.
The four copper atoms in 60 form a ‘‘butterfly’’ structure bridged by two ditelluride ligands (TeaTe 2.800–2.812 Å). Every copper atom is additionally coordinated by a phosphine ligand.
Clusters 61 and 62 are spherical molecules based on tellurium polyhedra, parts
of which are strongly distorted because of the presence of TeaTe bonds. Both
compounds display Frank-Kasper polyhedra containing 15 (61) or 16 (62) tellurium
atoms centered by one additional tellurium ligand. The tellurium polyhedron of 61
is similar to that observed in 20, although it differs in the presence of three TeaTe
bonds (2.936–3.011 Å). The copper atoms are bonded to the tellurium atoms in a
different arrangement to that in 20. Eight [CuPtBu3 ] units are situated above Te3
faces of the polyhedron, while two such units are edge-bridging. Eleven further
copper atoms are coordinated by three tellurium atoms, eight of which are localized below Te3 faces, while the other three bind between the central tellurium
151
152
3 Syntheses and Characterizations
59
61
60
62
Molecular structures of stoichiometric copper telluride clusters [Cu 4 Te(Te2 )2 (PCyPh2 )4 ] (59),
[Cu 4 (Te2 )2 (PiPr)4 ] (60), [Cu26 Te10 (Te2 )3 (PtBu3 )10 ] (61), and
[Cu28 Te13 (Te2 )2 (PEt2 Ph)12 ] (62) containing TeaTe bridges.
C and H atoms are omitted for clarity.
Fig. 3-59
atom and the tellurium atoms of the polyhedron. The remaining copper atoms are
localized in the center of the molecule and are each tetrahedrally surrounded by
four tellurium atoms.
In 62, two of the copper centers display relatively high thermal parameters, from
which it can be concluded that their positions are only partly occupied. A deficiency of electrons at the tellurium atoms resulting from the partial occupation of
copper sites could either be formally delocalized in the valence band or localized as
a ditelluride group following a structural deformation similar to a Peierls deformation [64]. Considering the tellurium polyhedra, the molecular structure of 62
shows two short TeaTe bonds (2.848 and 2.876 Å) and four TeaTe distances in the
range 3.150–3.351 Å. The latter tellurium atoms have thermal ellipsoids which are
elongated in the direction of the formal TeaTe bond. As TeaTe bonds and copper
centers seem to be disordered to a certain extent in 62, the given formula represents the idealized composition of a stoichiometric cluster compound with three
TeaTe bonds. The situation in 62 can therefore be viewed as being somewhat sim-
3.1 Semiconductor Nanoparticles
ilar to that in CuS, a well-known mixed-valence compound that can be described as
(Cuþ )2 Cu 2þ (S2 ) 2 S 2 or as (Cuþ )3 (S2 ) 2 S [65].
In a different approach to the synthesis of copper telluride cluster compounds,
we used alkyl- and aryl-(trimethylsilyl)tellurium compounds RTeSiMe3 (R ¼
organic group). Because of the facile cleavage of the carbon-tellurium bond, one
obtains clusters featuring either exclusively TeR ligands, a mixture of TeR and
Te 2 ligands, or solely Te 2 ligands. The tellurium atoms of the tellurolato ligands
bridge two or three copper atoms at most, whereas ‘‘naked’’ Te 2 ligands can coordinate to up to seven copper atoms. The polyhedra of tellurium atoms usually
possess unusual structures, even in the smaller clusters. In the larger clusters one
already observes characteristics of layer type arrangements of tellurium atoms. As
in the cluster compounds discussed earlier, the copper atoms are coordinated in
linear, trigonal or tetrahedral geometries by tellurium atoms. Metal centers coordinated by two or three chalcogens may also be ligated by a phosphine ligand. The
molecular structures of the copper telluride/tellurolate clusters 63, 64, 66, 71 and
72 are given in Figure 3-60.
The ionic cluster compound 63 contains six TePh ligands, the tellurium atoms
of which define an octahedron. In contrast to 64 (see below), all tellurolato ligands
act as m2 -bridges. There are two distinct coordination geometries around the copper. Two of the metal centers are located inside the cluster below Te3 faces. The
others are coordinated by three tellurium atoms above Te3 faces and are also each
ligated by one phosphine ligand each.
The structure of 64 is also based on a non-bonded octahedral array of tellurolato
ligands. With the exception of one tellurium atom which forms two CuaTe bonds,
all of these are bridging three copper sites. Four of the copper atoms adopt tetrahedral geometries, either with three tellurium atoms and one phosphorus or
through two CuaTe and two CuaP bonds. The other two copper atoms lie in opposite deltahedral Te3 faces of the Te6 octahedron. 64 is therefore related to the recently published homoleptic hexanuclear copper selenolate complex [Cu 6 {Se(2,4,6iPr3 C6 H2 )}6 ] [8].
Clusters 66 and 67 are identical except for the organic groups R of the PR3
ligands. Both contain two tellurolato layers of three TePh groups each and one
central tellurido layer of three Te 2 ligands. The resulting Te9 polyhedron can
alternatively be viewed as two face-sharing Te6 octahedra. Six of the twelve copper
atoms are coordinated in a distorted trigonal fashion by three tellurium atoms,
resulting in shorter contacts to Te 2 ligands than to the TePh . The other six
copper atoms are tetrahedrally surrounded by three tellurium atoms and one
phosphine ligand. Compounds 66 and 67 may be viewed as condensation products
of two (TePh)6 frameworks like those found in 64.
The synthesis of the ionic cluster compound 71 is the only example to date of a
co-condensation reaction of two smaller clusters, namely 67 and 63, to form a
larger one. However, the structure of 71 cannot be considered as being built up
from recognizable structural fragments of the precursor compound. The tellurium
atoms form an unusual polyhedron consisting of twelve TePh ligands and nine
Te 2 ligands. Two of the tellurolato ligands act as a m4 -bridges, while the other ten
153
154
64
71
Molecular structures of
stoichiometric copper telluride/tellurolate
clusters [Cu5 (TePh)6 (PEt)3 ][PEt3 Ph] (63),
[Cu 6 (TePh)6 (PEtPh2 )5 ] (64),
[Cu12 Te3 (TePh)6 (PPh3 )6 ] (66),
Fig. 3-60
66
72
[Cu29 (TePh)12 Te9 (PEt3 )8 ][PEt3 Ph] (71), and
[Cu50 Te17 (TePh)20 (PEtPh2 )8 ][PEtPh3 ]4 (72)
containing TeR ligands. C and H atoms are
omitted for clarity.
3 Syntheses and Characterizations
63
3.1 Semiconductor Nanoparticles
act as m3 -bridging atoms between the copper centers. The telluride ligands each
coordinate six, seven, eight, or ten copper atoms with CuaTe distances between
2.58 and 3.14 Å. Eight of the copper atoms which are bonded to phosphine ligands
cap Te3 faces at the cluster surface, while four further copper atoms without phosphine ligands are coordinating to Te3 faces from the inside of the cluster core. The
remaining seventeen copper atoms are surrounded in a tetrahedral manner in the
center of the molecule. Interestingly, a structurally related silver tellurium compound, [Ag30 (TePh)12 Te9 (PEt3 )12 ] [71] can be prepared by the direct reaction of
AgCl with Te(SiMe3 )2 /PhTeSiMe3 , which is in contrast to the synthesis of 71. This
cluster represents the only example of a copper-telluride-tellurolate cluster yet isolated with a distinct structural relationship with its silver analog.
Upon standing in daylight, solutions of 64 in benzene gradually darken to a
brown color. Over the course of several days, brown platelets of the ionic, mixed
tellurido/tellurolato cluster 72 crystallize from these solutions. On the other hand,
when solutions of 64 in benzene are protected from light, they show no sign of
darkening even after several weeks at room temperature. The detection of a
125
Te{ 1 H}-NMR-signal at d ¼ 688 ppm reveals the formation of TePh2 in the solution, giving good evidence for the mechanism of formation of the naked tellurido
ligands. The most striking feature of the cluster core is the manner in which the
tellurium ligands are arranged: the molecule consists of ‘‘top’’ and ‘‘bottom’’ layers
of m3 -TePh ligands with a central tellurido (Te 2 ) layer. Sandwiched between
these are the 50 copper atoms along with the remaining seven tellurido ligands.
The 20 phenyltellurolato ligands are each bonded to three copper atoms and, in
conjunction with the eight coordinated PEt2 Ph ligands, serve to effectively stabilize
the inner copper telluride core. The telluride ligands of the middle Te10 layer act as
m6 - and m7 -bridges between the copper sites. The seven remaining Te 2 centers
function as m6 -, m7 - or m8-ligands. Eight of the copper atoms are bonded to Te3 faces
at the surface of the cluster molecule and to one phosphine ligand. The 42 copper
atoms in the center of the molecule are either trigonally (18) or tetrahedrally (24)
coordinated by the tellurium neighbors.
Counting the electrons for 73–85 on the assumption of Te 2 and Cuþ leads to an
overall electron excess for 79 and an electron deficiency for the other compounds.
These compounds can therefore be thought of as being of mixed-valence type. The
molecular structures of 73, 76, 79, 83, and 85 are depicted in Figure 3-61.
The smallest mixed-valence cluster molecules, 73–75, feature 16 copper atoms
arranged around a Te9 polyhedron. However, the tellurium substructure cannot be
derived from the skeleton of 57 by simply adding one tellurium atom, as one
might have assumed. Two of the copper atoms that carry a phosphine ligand
bridge TeaTe non-bonding edges. The other six [CuPR2 R 0 ] units bridge three tellurium atoms, the copper atoms thus being tetrahedrally coordinated. The remaining eight copper atoms are coordinated by three tellurium atoms in a distorted trigonal planar manner. Assuming that the tellurium atoms have a formal
charge of 2, one can formally assign fourteen Cu 1þ centres and two Cu 2þ . However, the co-existence of Cu 2þ with Te 2 is unlikely for thermodynamical reasons
because of an electron transfer process from the reducing Te 2 ligand to the oxi-
155
156
76
83
Molecular structures of mixed-valence copper
telluride clusters [Cu16 Te9 (PPh3 )8 ] (73), [Cu23 Te13 (PPh3 )10 ] (76),
[Cu26 Te12 (PEt2 Ph)12 ] (79), [Cu44 Te23 (PPhnPr2 )15 ] (83), and
[Cu58 Te32 (PCy3 )16 ] (85). C and H atoms are omitted for clarity.
Fig. 3-61
79
85
3 Syntheses and Characterizations
73
3.1 Semiconductor Nanoparticles
dizing Cu 2þ center. An assignment of 16 Cuþ , seven Te 2 and two Te 1 is not
supported by structural means. The compounds should therefore rather be described in terms of a ‘‘pure’’ Cuþ substructure, ligated by tellurium centers with an
average charge of 1.78. This electron deficiency can be delocalized in the valence
band of the compound, as has recently been shown by calculations on the electronrich compound 79 (see below) [66].
Thirteen tellurium atoms form a distorted, centered icosahedron in the closely
related compounds 76–78. Ten of the 20 Te3 faces are bridged by copper-phosphine
units, whereas ten more copper atoms are situated below the remaining faces, either trigonally coordinated or tetrahedrally surrounded, because of an additional
CuaTe bond to the central tellurium atom. Compound 78 needs twelve phosphine
ligands for a sufficient shielding of the cluster core. This can be put down to the
smaller steric demand of the phosphine PEt3 compared to PPh3 . As a result, in
78, twelve of the copper atoms are coordinated above Te3 faces and eight below.
The three remaining copper atoms in all three compounds are positioned inside
the icosahedron. They each have three tellurium neighbors, two of which belong
to the Te12 shell, while one is the central tellurium ligand. Formal counting of
charges leads to an electron excess of three, and, assigning a charge of þ1 to every
copper atom as discussed above, an average charge of 1.77 is obtained for the 13
tellurium ligands.
An icosahedron of tellurium atoms is also found in 79, but, in contrast to those
observed in 76–78, it does not contain a central atom. The arrangement of the
copper atoms above the Te3 faces is similar to that found in 78. This might have
been anticipated, given the Tolman angle for both phosphines. Twelve [CuPEt2 Ph]
units are bonded above Te3 faces, and eight copper atoms cap the other Te3 faces
from the interior. In contrast to the situation in 76–78, in 79 six rather than three
copper atoms are located in the center of the cluster molecule forming an octahedron with intermetallic distances between 2.592 and 2.607 Å. These distances lie
within the range of CuaCu distances measured in the other copper telluride clusters under investigation. Therefore, it cannot be determined from these distances
whether d 10 –d 10 interactions are present within the Cu 6 octahedron. However, it is
remarkable that two of the copper atoms must be formally uncharged Cu 0 atoms
lying alongside 24 Cuþ centers, if one assigns a formal charge of 2 to all the tellurium ligands. Calculations of the electronic band structure reveal that a localization of the oxidation states does not occur and that the excess electron density
cannot be assigned to d 10 –d 10 interactions [66]. The ‘‘additional’’ electron pair,
which arises if one assigns exclusively Cuþ and Te 2 centers, is better described
as an MO embedded in the valence band of tellurium 5p orbitals and copper 4s
orbitals.
The largest copper telluride cluster molecules synthesized to date are also mixedvalence compounds. Compounds 82–87 display ellipsoidal tellurium deltahedra
with an increasing number of inner tellurium atoms, ranging from two inner tellurium atoms in 82–84 to four such tellurium atoms in the centers of the molecular structures of 85–87. As usual, the copper atoms that are coordinated by phosphine ligands are coordinated above the Te3 triangles at the cluster surface. The
157
158
3 Syntheses and Characterizations
other copper atoms are either coordinated by three or, especially in the cluster
center, by four tellurium ligands. The structures of 85–87 already suggest the formation of distorted layers of tellurium atoms showing incipient characteristics of
bulk Cu2 Te [40]. However, the turning point has not yet been reached for copper
telluride cluster compounds. In view of the reported tendency for ever bigger molecules to be necessary in order that bulk structure properties will be observed on
going from CuaS to CuaSe clusters, the cluster size that might begin to show bulk
Cu2 Te structural characteristics should in fact be considerably larger than a ‘‘Cu58’’
species and probably even larger than a ‘‘Cu70’’ analog of the turning point to
‘‘large’’ clusters found for the copper selenide system.
Chalcogen-Bridged Silver Clusters
The development of the chemistry of Cu2 S, Cu2 Se and Cu2 Te complexes has yet to
be mirrored for the heavier metals Ag and Au. Monodentate phosphine ligands
have generally been shown to be incapable of providing a kinetically stabilizing
sphere when Ag(I) and Au(I) salts are reacted with E(SiMe3 )2 , and these reactions
lead instead to the formation of amorphous binary solids.
It is clear from the development of the chemistry of the copper chalcogenide
cluster complexes described above that (i) the incorporation of additional surface
ligands in the form of SeR and TeR functionalities and/or (ii) the use of bidentate
phosphine ligands offer additional structural flexibility. Additionally, they may be
used to stabilize AgSe, AgTe and AuSe cluster cores, themselves synthesized with
bis(silylated) reagents. The use of monosilylated chalcogen reagents, as the hydrocarbon moiety is relatively inert and unreactive toward metal salts, could thus aid
in inhibiting the formation of the bulk, amorphous solids. Furthermore, the reaction chemistry of metal-selenolate and -tellurolate complexes can also provide the
source of E 2 ligands required for the formation of larger polymetallic complexes
[19, 61]. Thus, the use of monsilylated reagents RESiMe3 (R ¼ aryl, alkyl) has led
to the development of nanometer scale Ag/Se/SeR and Ag/Te/TeR complexes. The
carbon chain on the chalcogen center is readily modified, and, as discussed below,
this plays an important role in governing the structures formed.
3.1.3.3
Silver-Selenium Clusters
The reaction chemistry of Ag(I) and RESiMe3 complexes in the presence of tertiary
phosphine ligands is summarized in Scheme 3-9. Of note is the requirement for
acetate salts of silver for the successful isolation of polynuclear complexes, as AgCl
invariably leads to Ag2 Se formation. This contrasts with the reaction chemistry of
Cu(I) described above and also with the chemistry of silver-tellurium complexes
(see below). Of equal importance is the facile formation of selenide (Se 2 ) ligands
from RSe fragments, as this allows for the formation of extremely large AgaSe
frameworks. Although this reactivity limits the number of low nuclearity complexes that can be isolated, and thus limits the amount of information that can be
obtained about condensation pathways, it offers the distinct advantage of the formation of the largest structurally characterized metal-chalcogenide complexes reported to date.
3.1 Semiconductor Nanoparticles
−40 °C, C5H12
R = Ph
R’ = Pr
R” = tBu
AgO2CR
+ R”SeSiMe 3
+
PR’3
– SiMe3O2CR
[Ag30Se8(SetBu)14(PnPr3)8]
88 [67]
[Ag90Se38(SetBu)14(PEt3)22]
89 [67]
[Ag114Se34(SenBu)46(PtBu3)14]
90 [67]
[Ag112Se32(SenBu)48(PtBu3)12]
91 [67]
−40 °C, THF
R = Ph
R’ = Et
R” = tBu
−30 °C, C5H12
R = C11H23
R’ = tBu
R” = nBu
−30 °C,
C5H12/THF
R = Me
R’ = tBu
R” = nBu
Scheme 3-9 The synthesis of selenido/selenolato-bridged silver
clusters with monodentate PR3 ligands.
88
Fig. 3-62
The molecular structure of [Ag30 Se8 (SetBu)14 (PPr3 )8 ] (88).
The reaction of silver benzoate with tBuSeSiMe3 and PPr3 at 40 C in pentane
solvent affords the light-sensitive, spherical cluster [Ag30 Se8 (SetBu)14 (PPr3 )8 ] 88 in
good (60%) yield [67]. The structure of 88 (Figure 3-62) illustrates the surface
groups (PR3 and tBu) stabilizing the AgSe core. The SetBu ligands bridge two and
three silver centers and the Se 2 ligands adopt coordination numbers five and six.
159
160
3 Syntheses and Characterizations
89
Fig. 3-63
The core atoms in [Ag90 Se38 (SetBu)14 (PEt3 )22 ] (89).
There are ten silver centers that are two-coordinate, bonded to two SetBu ligands,
two Se 2 ligands, or one SetBu ligand and one Se 2 ligand. The remaining twenty
silver atoms, including the eight that are bonded to PPr3 ligands, exhibit distorted
trigonal planar or tetrahedral coordination geometries. As is the case for the stoichiometric copper-chalcogenide complexes described in the preceding sections, the
shortest non-bonding contacts (2.904–2.991 Å) between the metal centers are consistent with the þ1 oxidation states assigned to all Ag atoms.
A similar reaction for the formation of 88 with PEt3 in THF leads to the larger
cluster [Ag90 Se38 (SetBu)14 (PEt3 )22 ] (89) in 60% yield (Figure 3-63) [67]. The similar
basicities and Tollman cone angles [42] for PPr3 and PEt3 suggest that the use of a
different phosphine ligand for 89 does not account for the different product formation. A more likely reason is the combination of a different solvent (THF vs
pentane) with higher reaction/crystallization temperatures (20 C for 89; 60 C
for 88). Higher reaction temperatures for these clusters lead to the formation
of bulk Ag2 Se and other unidentified amorphous materials. The thirty-eight Se 2
ligands in the core of 89 result in a larger, pseudo-spherical framework with core
dimensions of 1.1 1.7 1.9 nm 3 . The 90 Ag centers are ligated to the 38 Se 2
and 14 SetBu groups, with 22 also bonded to the surface PEt3 ligands. The metal
atoms display linear, trigonal planar, and tetrahedral coordination. The SetBu
groups ligate in m2 - and m3 -fashion, and the Se 2 centers also exhibit varying coordination modes, with those in the outer shell bridging four and five silver atoms
and the core ligands forming five to eight bonds to the metals. The arrangement of
the selenium centers in 89 is best described in two ways (Figure 3-64). The
projection at the top of Figure 3-64 illustrates the approximately D6h symmetry
found in the non-bonded Se52 polyhedron, forming two ‘‘belts’’ of selenium
ligands. The outer belt consists of 12 SetBu and 18 Se 2 , whereas the inner one is
3.1 Semiconductor Nanoparticles
Fig. 3-64
Two projections of the Se52 selenium sub-structure in 89.
composed exclusively of 18 Se 2 . A perpendicular view of the Se framework emphasizes the development of a layered arrangement of these ligands, which dominates the overall structural motif observed in the larger AgaSe cluster complexes
described below.
The 2.3 2.3 2.5 nm 3 AgSe core in [Ag114 Se34 (SenBu)46 (PtBu3 )14 ] 90 (Figure
3-65) is reproducibly formed from the reaction of silver laurate (C11 H23 CO2 Ag),
PtBu3 and nBuSeSiMe3 in pentane at 30 C [67]. The increased size of 90 compared with 88 and 89 is accompanied by a rather marked change in the gross
structural features exhibited. Specifically, the arrangement of Se atoms in 90
(Figure 3-66) corresponds to a distorted cubic-centered structure that is also observed in solid Ag2 Se. The Ag atoms have trigonal and tetrahedral coordination
and are distributed between the layers of Se, which is also consistent with the
structural features observed in the bulk solid. This ‘‘transition’’, i.e. the sizedependent structural changes, is also observed in the PR3 -stabilized Cu2 Se clusters, where a clear relationship between molecule and bulk is observed for structures containing 70 or more Cu atoms, as described in the preceding sections. One
cannot ignore, however, the effect of the different substituents on the selenolate
ligands (nBu vs tBu), which cannot be discounted. The nBuSe ligands in 90
bridge three and four Ag(I) centers, and the selenide ligands again are within the
bonding distance range to 5, 6, 7 and 8 metals.
The cluster [Ag112 Se32 (SenBu)48 (PtBu3 )12 ] 91 (Figure 3-65) is closely related to
that of 90, and is formed from the reaction of AgOAc, PtBu3 and nBuSeSiMe3 in
pentane/THF mixtures at 30 C [67]. The number of selenium atoms in 90 and
161
162
3 Syntheses and Characterizations
90
91
Fig. 3-65 The core structures in [Ag114 Se34 (SenBu)46 (PtBu3 )14 ]
(90) (top) and [Ag112 Se32 (SenBu)48 (PtBu3 )12 ] (91) (bottom).
91 is identical; however, the number of surface selenolate groups is greater in the
latter (48 vs 46), with a concomitant reduction in the number of PtBu3 ligands (12
vs 14). Two fewer (112) Ag centers in 91 relative to 90 is consistent with the lower
Se 2 :nBuSe ratio in the former (32:48 and 34:46, respectively).
Bidentate phosphine ligands have been successfully employed for the synthesis
of Cu/E/ER complexes (E ¼ Se, Te) from silylated reagents, as described above.
Their utility in stabilizing Ag/SeR and Ag/Se/SeR frameworks using analogous
reaction strategies is demonstrated and summarized in Scheme 3-10.
The use of bidentate phosphine ligands in conjunction with low temperature
conditions is apparently sufficient in enabling the isolation of low-nuclearity Ag/
3.1 Semiconductor Nanoparticles
Fig. 3-66
The selenium frameworks in 90 (top) and 91 (bottom).
SeR complexes from Ag(I) salts and RESiMe3 , hitherto impossible when working
with PR3 ancillary ligands. For example, the cluster [Ag 4 (SeiPr)4 (Ph2 PCH2 PPh2 )2 ]
(92) is isolated in 45% yield from silver thiocyanate (or silver benzoate), bis(diphenylphosphino)methane and iPrSeSiMe3 at low temperatures [68]. Cluster 92
(Figure 3-67) consists of an Ag 4 tetrahedron with each non-bonded Ag Ag edge
bridged by a m2 -iPrSe ligand. The trigonal planar coordination geometry about the
Ag(I) centers is completed via coordination to a phosphine ligand.
The m3 -bridging capacity of alkylselenolate ligands and the flexible coordination ability of bis(diphenylphosphino)propane ligands lead to an infinite repeat of
[Ag8 (SeEt)8 (Ph2 PCH2 CH2 CH2 PPh2 )] units (Figure 3-68) in 93 [68]. This poly-
163
164
3 Syntheses and Characterizations
[Ag4(SeiPr)4(dppm)]2
92 [68]
[Ag8(SeEt)8(dppp)]∞
93 [68]
[Ag124Se57(SePtBu2)4Cl6(tBu2P(CH2)3PtBu2)12]
94 [68]
[Ag28Se6(SenBu)16(dppp)4]
95 [68]
[Ag172Se40(SenBu)92(dppp)4]
96 [67]
−60 °C→ −20 °C,
Et2O
X = SCN
R = iPr
P∩ P =
Ph2PCH2PPh2
AgX
+
P∩ P
+ RSeSiMe3
– SiMe3X
−40 °C→ RT,
C7H8/COD
X = O2CPh
R = Et
P∩ P =
Ph2P(CH2)3PPh2
−20 °C→ RT,
CH2Cl2
X = O2CC11H23
R = tBu
P∩ P =
tBu 2P(CH2)3PtBu2
−80 °C→ −20 °C,
CH2Cl2
X = O2CC11H23
R = nBu
P∩ P =
Ph2P(CH2)3PPh2
−20 °C→ RT,
CHCl3/C7H8
X = O2CC11H23
R = nBu
P∩ P =
Ph2P(CH2)3PPh2
The synthesis of selenido/selenolato-bridged
silver clusters with bidentate phosphine ligands.
Scheme 3-10
meric structure consists of a strand of distorted Se6 octahedral arrangements of
ethylselenolate ligands, each of the EtSe ligands functioning as m3 -bridges to the
Ag(I) centers. There is no discrepancy between ‘‘intra’’ and ‘‘inter’’ unit bonding
distances, which range from 2.509(4) to 2.847(3) Å. In contrast to Ag 4 cluster 92,
the silver atoms in the Ag8 cluster 93 display different coordination geometries: a
linearly coordinated Ag is bonded to two EtSe ligands above an edge and a trigonal
planar Ag center resides in a deltahedral face of the Se6 octahedron; two additional
silver atoms show trigonal planar geometry ‘‘between’’ the octahedral subunits; the
remaining silver atoms are three- and four-coordinate, occupying positions above
and within the Se6 octahedra. Although the ethyl groups bonded to the selenium
sites have poor shielding ability, the AgSe strand is effectively insulated and stabilized by the DPPP ligands (Figure 3-69).
3.1 Semiconductor Nanoparticles
Fig. 3-67
92
The molecular structure of [Ag 4 (SeiPr)4 (Ph2 PCH2 PPh2 )2 ] (92).
Fig. 3-68
A projection of the infinite repeat in [Ag8 (SeEt)8 (Ph2 PCH2 CH2 CH2 PPh2 )]y (93).
Fig. 3-69
A space-filling model of polymeric 93.
165
166
3 Syntheses and Characterizations
94
A projection of the molecular structure of
[Ag124 Se57 (SePtBu2 )4 Cl6 (tBu2 P(CH2 )3 PtBu2 )12 ] (94) (carbon
atoms omitted).
Fig. 3-70
Much larger AgaSe frameworks are formed using bidentate ligands if reaction
conditions and the selenium reagent are modified. The cluster [Ag124 Se57 (SePtBu2 )4 Cl6 (tBu2 P(CH2 )3 PtBu2 )12 ] (94, Figure 3-70) is isolated as black crystals
from the reaction of silver laurate, 1,3-bis(ditertbutylphosphino)propane, and tertbutylselenotrimethylsilane in methylene chloride at 20 C [68]. Cluster 94, with
a spherical @2 nm AgSe core, is a very unusual member of the series of highnuclearity silver selenide clusters, bears no structural relationship to the clusters
90 and 91 described above, and displays no structural characteristics of the bulk
phases of Ag2 Se.
The important structural features of 94 are the Frank-Kasper polyhedra surrounding the central selenide ligand. A non-bonded (Se Se ¼ 4.93–5.36 Å) Se16
polyhedron surrounds the central Se 2 and is itself enveloped by 44 selenium
centres, four of which are from the surface SePtBu2 ligands (Figure 3-71). This
defined Se44 polyhedron (Se Se ¼ 3.80–5.99 Å) possesses 84 D3 faces that can
be described as a growth of the inner Se16 deltahedron. The Se 2 ligands in 94 act
as m3 , m4 , m5 , m6 , m7 or m8 bridges to the silver atoms, with the AgaSe distances
ranging from 2.491(3) to 3.101(3) Å. The four selenoditertbutylphosphido ligands,
SePtBu2 , are m3 -h 1 (P):h 2 (Se) bonded to 12 Ag(I) centers. The formation of these
ligands from the reaction conditions must result from the decomposition of the
tBuSeSiMe3 and tBu2 P(CH2 )3 PtBu2 ligands. For their part, the six surface chlorides, which arise from the reaction solvent methylene chloride, are m3 and m2
bridging to the silver centers.
3.1 Semiconductor Nanoparticles
Fig. 3-71 The selenium substructure in 94. The central Se is
surrounded by the Se16 Frank-Kasper polyhedron, which is itself
surrounded by the Se44 polyhedron in a second cooordination
sphere.
The combination of 57 Se 2 , 6 Cl and 4 SePtBu2 ligands again allows for
þ1 oxidation states to be assigned to the silver atoms. The 124 Ag(I) centers are
distributed as follows: twelve of the 28 tetrahedral sites of the central Se16 deltahedron surrounding the central Se are occupied by silver atoms in distorted geometries; 80 silver atoms located between the Se16 and Se44 polyhedra exhibit
linear, trigonal planar and tetrahedral coordination, and the remaining 32 silver
atoms are located at the cluster surface with coordination numbers three and four.
The effects of solvent and reaction/crystallization conditions on product formation are clearly evident in the syntheses of [Ag28 Se6 (SenBu)16 (dppp)4 ] (95) and
[Ag172 Se40 (SenBu)92 (dppp)4 ] (96). The former is obtained in 45% yield from
1:0.25:1 mixtures of AgO2 CC11 H23 :Ph2 P(CH2 )3 PPh2 :nBuSeSiMe3 in CH2 Cl2 at
80 C, slowly rising to 20 C [68]. In contrast, the same combination of reagents in CHCl3 /toluene at 20 C generates cluster 96 in 20% yield when reaction solutions are allowed to warm to room temperature [67].
The red crystals of 95 formed from the reaction mixtures are very sensitive and
decompose at temperatures above 10 C. X-ray analysis on the highly solvated
crystals is possible, however, and the structure of 95 (Figure 3-72) displays the features expected for a ‘‘medium-sized’’ AgaSe complex and is not structurally dissimilar to [Ag30 Se8 (SetBu)14 (PPr3 )8 ] (88). Each of the nBuSe ligands is m3 -bonded
on the surface of the cluster, and the six Se 2 atoms are arranged to form a nonbonded octahedron at the center of the cluster. For their part, the Ag(I) atoms display the (distorted) linear, trigonal planar and tetrahedral coordination geometries
expected in such a medium-sized cluster.
The cluster 2.4 2.4 2.0 nm 3 AgSe core in [Ag172 Se40 (SenBu)92 (dppp)4 ] 96
represents the largest structurally characterized metal-selenide complex reported to
167
168
3 Syntheses and Characterizations
Fig. 3-72
95
A projection of the core atoms in [Ag28 Se6 (SenBu)16 (dppp)4 ] (95).
96
Fig. 3-73 The Ag, Se and P atoms in the 2.4 2.4 2.0 nm 3
AgSe core in [Ag172 Se40 (SenBu)92 (dppp)4 ] (96).
date (Figure 3-73) [67]. With increasing cluster size, the distribution of the silver centers appears to become chaotic, and the refinement of the X-ray data for 96 is correspondingly problematic. In this respect, 96 reflects the structure of solid Ag2 Se,
where the Agþ cations are also disordered. On the cluster surface, 92 nBuSe ligands are m3 - and m4 -bonded to the silver atoms (SeaAg ¼ 2.498(3)–3.064(3) Å).
Noteworthy is the fact that there are only eight phosphine-silver bonds (four DPPP
ligands) on the surface, as the butyl chains of the selenolate ligands provide the
majority of the stabilization of the silver-selenide core. The 40 Se 2 ligands act as
m5 , m6 , m7 and m8 bridges (AgaSe ¼ 2.44–3.30 Å) and are arranged with the Se centers of the nBuSe groups to form layered arrangements of selenium.
3.1 Semiconductor Nanoparticles
AgCl
+ nBuTeSiMe3
+
– SiMe3Cl
x PEt3
[Ag6(TenBu)6(PEt3)6]
97 [69]
[Ag32Te7(TenBu)18(PEt3)6]
98a [69]
−40 °C→ −30 °C,
C5H12
x=3
−40 °C→ −30 °C,
C5H12
x=1
[Ag48Te12(TenBu)24(PEt3)14]
99 [69]
−40 °C→ −30 °C,
C5H12
x=2
Scheme 3-11
A summary of the reaction chemistry of Et3 P/AgCl and nBuTeSiMe3 .
97
Fig. 3-74
The molecular structure of [Ag6 (TenBu)6 (PEt3 )4 ] (97).
Silver-Tellurium Clusters
Silver chloride is dissolved readily in ether or hydrocarbon solvents using excess
trialkylphosphine ligands. At room temperature, the addition of nBuTeSiMe3 to
these solutions proceeds with the precipitation of insoluble, brick-red, amorphous
materials or the formation of Ag2 Te. Low-temperature reactions again provide a
route to access crystalline products suitable for X-ray analysis (Scheme 3-11)
[69]. The addition of nBuTeSiMe3 to pentane solutions of a 1:3 mixture of
AgCl:PEt3 at 40 C yields a gold-colored solution from which bright yellow crystals of [Ag6 (TenBu)6 (PEt3 )4 ] (97, Figure 3-74) grow within several days in modest
yield (30%). The six tellurolate ligands define a non-bonded octahedron, with four
of the nBuTe ligands m3 -bonded and two nBuTe ligands bridging two Ag atoms.
Two Ag(I) centers are tetrahedral, bonded to two PEt3 and two nBuTe ligands, and
the four remaining silver atoms are trigonal planar and bonded to the butyltellurolate ligands only, residing in the deltahedral faces of the Te6 polyhedron. The
169
170
3 Syntheses and Characterizations
98a
Fig. 3-75
The molecular structure of [Ag32 (TenBu)18 Te7 (PEt3 )6 ] (98a) (carbon atoms omitted).
importance of phosphine:metal ratio in the synthesis of metal-chalcogenide
complexes with silylated reagents is again highlighted with the synthesis of
[Ag32 (TenBu)18 Te7 (PEt3 )6 ] (98a): if the ratio of phosphine to silver is only 1:1, there
is no evidence of 97, and cluster 98a is formed in low yield. The projection illustrates the layered nature of cluster 98a, which resides on a crystallographic inversion centre (Figure 3-75). The ‘‘top’’ layer of 98a consists of three m3 -nTe(Bu)
groups, whereas the second tellurium-rich layer contains three m3 -tellurolate and
three m5 -telluride (Te 2 ) ligands. The latter form a nearly planar hexagonal Te3 Ag3
ring (max. dev. ¼ 0.10 Å), which lies above (0.61 Å) a tetrahedral Ag center. The
central tellurium layer contains six m3 -nTe(Bu) ligands and one, central Te 2
atom, which is within the range of bonding distances to eight silver atoms. The
arrangement of Te is repeated in the bottom half of the structure. The result is that
24 Ag atoms all form three bonds to neighbouring Te atoms and six Ag atoms
achieve trigonal planar coordination by bonding to two tellurolate ligands and one
phosphine ligand. The butyltellurolate ligands in 98 symmetrically bridge three
metal centers. The combination of eighteen tellurolate (18 ) and seven telluride
ligands (14 ) again suggests a þ1 oxidation state for the thirty-two silver atoms.
The yield of cluster 98a is low (@5%), although the synthesis is reproducible. There
appears to be some inherent stability associated with the AgaTe framework, as
analogous [Ag32 (TenBu)18 Te7 (PR3 )6 ] clusters can be isolated using other tertiary
phosphine ligands, and [Ag32 (TenBu)18 Te7 (PEt2 Ph)6 ] (98b), [Ag32 (TenBu)18 Te7 (PnPr3 )6 ] (98c) and [Ag32 (TenBu)18 Te7 (PiPr3 )6 ] (98d) are also similarly prepared.
Unlike 98a–98c, for which yields are low, the yield of cluster 98d (R ¼ iPr) is rela-
3.1 Semiconductor Nanoparticles
Fig. 3-76
99
The molecular structure of [Ag 48 (TenBu)24 Te12 (PEt3 )14 ] (99) (carbon atoms omitted).
tively good (40%), as reaction solutions (of AgCl, PiPr3 and nBuTeSiMe3 ) can be
stirred at room temperature prior to crystallization without the formation of insoluble, solid materials. Presumably this facilitates the formation of additional Te 2
ions via tellurium-carbon bond cleavage of butyltellurolate moieties. The range of
cone angles of the phosphines 132–160 for which identical cluster frameworks are
observed illustrates the prominent role the alkyl chains on the Te centers play in
stabilizing the cluster core, consistent with the limited number (six) of phosphines
on the surface of the cluster. Attempts to improve the yields of clusters 98a–d by
using a mixture of nBuTeSiMe3 and Te(SiMe3 )2 does not yield any crystalline materials.
Modifying the ratio of triethylphosphine to silver chloride such that AgCl:PEt3 :
nBuTeSiMe3 is 1:2:1, the cluster [Ag 48 (TenBu)24 Te12 (PEt3 )14 ] (99, Figure 3-76) is
isolated as dark red crystals after several days at low temperatures. The cluster
consists of two spherical Ag24 Te16 sub-units linked by four m3-TenBu ligands
about a crystallographic inversion center. In contrast to 98, the fourteen AgaPEt3
centres in 99 are all tetrahedral. The remaining silver atoms within the cluster core
have coordination number three (trigonal planar) or four (tetrahedral), and longer
Ag Te contacts also suggest intermediate bonding descriptions. The 24 nBuTe
ligands and 12 Te 2 atoms are in agreement with a þ1 oxidation state to be assigned
to the Ag sites, and the Ag(I) Ag(I) contacts range between 2.802(2) and 3.222(2) Å.
Although the synthesis of butyltellurolate-silver complexes requires low temperatures to avoid the precipitation of amorphous solids and we have isolated relatively few structural types of these compounds (see above), the chemistry of phenyltellurolate-silver complexes from PhTeSiMe3 and AgCl is rich and varied with
respect to the polynuclear frameworks formed. These aryltellurolate reagents can
171
172
3 Syntheses and Characterizations
[Ag10(TePh)10(PMe3)2]∞
RT, C5H12
R = R’ = Me
AgCl
100 [70]
+
[Ag14Te(TePh)12(PMe3)8]
101 [70]
[Ag14Te(TePh)12(PEt2Ph)8]
102 [70]
+ PhTeSiMe3
+
2 PR2R’
– SiMe 3Cl
+
RT, THF/C5H12
R = Et, R’ = Ph
[Ag9(TePh)9(PEt2Ph)6]
103 [70]
The synthesis of tellurido/tellurolato-bridged
silver clusters with monodentate PR3 ligands.
Scheme 3-12
also be classified as ‘‘gentler’’ in that reactions between the phosphine-solubilized
metal salt and the silylated chalcogens can often be performed at room temperature to yield medium-sized polynuclear complexes.
The reaction of a suspension of AgCl:2 PMe3 with PhTeSiMe3 yields two products, [Ag10 (m3 -TePh)10 (PMe3 )2 ]y (100) and [Ag14 Te(TePh)12 (PMe3 )8 ] (101) and in
40% and 20% yield respectively (Scheme 3-12) [70].
Complex 100 consists of a polymeric one-directional arrangement, in which the
PhTe ligands serve as both intra- and intermolecular bridges. A crystallographic
inversion center relates the two halves of the Ag10 cluster frame (Figure 3-81),
which is repeated to form the polymeric chain. All PhTe ligands are bonded to
three Ag(I) atoms, and the silver centers are three-coordinate and trigonal planar
(sum of the angles ¼ 350.3–357.2 ) with the exception of those bonded to the
PMe3 ligand, which adopt a tetrahedral coordination geometry. There is no significant difference in length between the inter- and intracluster TeaAg contacts. The
structure of 100 consists of a rather shell-like make-up: the inner core consists of
silver atoms surrounded by a Te/Ag layer which is itself surrounded by PMe3 and
Ph ligands (Figure 3-77). The structure consists of a repeating clustered arrangement of silver and tellurium forming an AgTe ‘‘wire’’.
The second product formed in the reaction of AgCl:2 PMe3 with PhTeSiMe3 is
the molecular complex [Ag14 Te(TePh)12 (PMe3 )8 ] (101) [70]. In 101 (Figure 3-78),
the fourteen silver centers are bonded to m3 -, three m4 - and two m-TePh ligands with a central m6 -Te 2 . The AgaTe distances associated with the two m-TePh
ligands are the shortest (on average) within the cluster. The non-bonded Te13
framework can be described as a capped, centered icosahedron in which one m5 vertex is missing. The eight silver centers bonded to phosphine ligands exhibit
tetrahedral coordination geometry, the remaining sites being either trigonal planar
or tetrahedral.
The highly distorted and asymmetric nature of the cluster 101 can be contrasted
with the framework observed in [Ag14 Te(TePh)12 (PEt2 Ph)8 ] (102, Figure 3-79). Although the ratio of Ag:TeR:Te:PR3 is the same in 101 and 102, there are marked
differences in the structures of the two clusters [70]. The eight PEt2 Ph ligands in
3.1 Semiconductor Nanoparticles
Fig. 3-77 Two different projections of the infinite repeat
observed in [Ag10 (m3 -TePh)10 (PMe3 )2 ]y (100).
Fig. 3-78
101
1
The AgaTeaP framework in [Ag14 Te(TePh)12 (PMe3 )8 ] (101).
173
174
3 Syntheses and Characterizations
Fig. 3-79
102
The AgaTeaP framework in [Ag14 Te(TePh)12 (PEt2 Ph)8 ] (102).
101 are located at the corners of a cube, and the twelve TePh ligands are evenly
distributed, forming three (distorted) layers and each bonded to three silver atoms.
The central telluride is within the range of bonding distances (2.768(1)–2.825(1) Å)
to six silver atoms. The Te13 framework can be described as a highly distorted centered icosahedron. Of the twenty deltahedral faces of the icosahedron, eight are
capped by the tetrahedral Ag(PEt2 Ph) sites. The AgaTe distances at these fourcoordinate silver sites are longer than the six three-coordinate Ag(I), which lie
within the Te13 frame: these latter have AgaTe contacts ranging from 2.660(1) to
2.825(1) Å, and for the former set the AgaTe contacts are longer, ranging from
2.838(1) to 3.026(1) Å. The coordination about the three-coordinate silver centers is,
however, highly asymmetric, as invariably one of the TeaAgaTe angles lies close
to 90 .
The second, and major, product (60%) from the reaction of AgCl:PEt2 Ph with
Te(Ph)SiMe3 is the cluster [Ag9 (TePh)9 (PEt2 Ph)6 ] (103a) (Scheme 3-12) [70]. The
AgaTe framework (Figure 3-80) sits on a 3-fold rotation axis that bisects three
three-coordinate Ag(I) atoms in the center of the rather columnar cluster that are
bonded to three TePh ligands in trigonal planar fashion. The other six silver
atoms are each bonded to one PEt2 Ph and three TePh ligands, and are tetrahedral. All TePh groups bridge three Ag(I) atoms. The coordination environment
about the three central TePh ligands is unusual in the AgaTePh clusters described, as it consists of ‘‘T-shaped’’ coordination about tellurium. The central layer
consists of a planar Ag(TePh)3 arrangement, with the three phenyl rings lying
in this plane. The extended packing arrangement of 103a results in the central
Ag(TePh)3 layers being aligned, in a mixed aromatic-AgaTe infinite repeat.
With the exception of [Ag32 (TenBu)18 Te7 (PiPr3 )6 ] (98d), systematically accessing
higher nuclearity, nanoscale AgTe clusters in good yield using monosilylated re-
3.1 Semiconductor Nanoparticles
Fig. 3-80
The AgaTeaP framework in [Ag9 (TePh)9 (PEt2 Ph)6 ] (103a).
agents R/ArTeSiMe3 is not as effective as the corresponding technique for CuaTe
or AgaSe complexes. Using a combination of silylated tellurium sources, however,
(i.e. ArTeSiMe3 and Te(SiMe3 )2 ) to access metal-telluride clusters in good yield can
be used if the substituent on the tellurolate ligand is an aryl group. Thus, although
the reaction of AgCl:PR3 solutions and Te(SiMe3 )2 leads to the formation of bulk
Ag2 Te, AgCl dissolved in ethereal solvents with triethylphosphine reacts with a
combination of Te(SiMe3 )2 and ArTeSiMe3 (Ar ¼ Ph, mes; mes ¼ C6 H2 Me3 ) to
yield the cluster complexes [Ag30 (TePh)12 Te9 (PEt3 )12 ] (104a) and [Ag 46 (Temes)12 Te17 (PEt3 )16 ] (105a), respectively (Scheme 3-13) [71].
When an AgCl:PEt3 (1:2) solution is treated with 0.5 equivalents of PhTeSiMe3
and 0.25 equivalents of Te(SiMe3 )2 at 40 C, red crystals of 104a form within a
few days at 30 C in fair yield (20%). The @1.3 1.1 0.8 nm 3 cluster core is
comprised of 12 tellurolate (PhTe ), 9 telluride (Te 2 ), and 30 silver atoms (Figure
3-81). Similar complexes can be prepared using the different phosphine ligands
(PEt2 Ph (104b), PnPr3 (104c)).
The order of addition of the silylated reagents does not affect the yield of product. The phenyltellurolate ligands in 104a act as both m3 - and m4 -bridges between
the Ag(I) centers. Each of the nine Te 2 atoms forms six bonding contacts to the
Ag centers (AgaTe ¼ 2.703(3)–3.185(3) Å). Of the twelve phosphine-bonded silver
sites, all but two, which are trigonal planar, exhibit tetrahedral coordination. The
coordination environments of the silver atoms bonded only to tellurium are varied:
two silver atoms form two bonds, each [J(TeaAgaTe) ¼ 153.1(1)–153.22(9) ], eight
silver atoms are three-coordinate and six are tetrahedral. This varied coordination
geometry about silver in the cluster contrasts with the trigonal planar geometry of
the majority of the silver centers in 98 and 99, but is consistent with the more
close-packed nature of the AgTe framework in 104 and the higher coordination
modes of the insterstitial Te ligands. Similar varied coordination geometries have
been observed for the copper-chalcogenide clusters described in the preceding sec-
175
176
3 Syntheses and Characterizations
− 40 °C → −30 °C
Et2O/THF
Ar = Ph
AgCl
+
2 PR2R’
[Ag30Te9(TePh)12(PEt3)12]
104a [71]
[Ag30Te9(TePh)12(PEt2Ph)12]
104b
[Ag30Te9(TePh)12(PnPr3)12]
104c
[Ag46Te17{Te(mes)}12(PEt3)16]
105a [71]
[Ag46Te17{Te(mes)}12(PEt2Ph)16]
105b
[Ag46Te17{Te(mes)}12(PnPr3)16]
105c
+ 0.5 ArTeSiMe3
+ 0.25 Te(SiMe3)2
– SiMe3Cl
−40 °C→ −30 °C
Et2O
Ar = mes
(mes = C6H2Me3)
The synthesis of higher nuclearity tellurido/
tellurolato-bridged silver clusters from AgCl, Te(SiMe3 )2 , and
ArTeSiMe3 .
Scheme 3-13
104a
Fig. 3-81
A projection of the core atoms in [Ag30 (TePh)12 Te9 (PEt3 )12 ] (104a).
tions: whereas for copper-sulfide clusters containing up to 28 Cu centres, near linear SaCuaS and trigonal planar S2 aCuaP arrangements are observed, the copper
centres in larger selenium clusters exhibit tetrahedral and trigonal planar coordination. Cluster 104a is structurally related to the copper-tellurium cluster [Cu29 (TePh)12 Te9 (PEt3 )8 ][PEt3 Ph] (71), which is formed from the co-condensation reaction of the two smaller clusters [Cu12 Te3 (TePh)6 (PEt3 )6 ] (63) and [Cu5 (TePh)6 3.1 Semiconductor Nanoparticles
105a
The molecular structure of [Ag 46 (Temes)12
Te17 (PEt3 )16 ] (105a) (carbon atoms omitted).
Fig. 3-82
(PEt3 )3 ][PEt3 Ph] (67) [61]. This is the only example identified to date where there is
a distinct structural relationship between a copper-telluride-tellurolate nanocluster
and a silver analog. There are four additional PEt3 ligands on the surface of 104a vs
63 and the phosphine-bonded silver centers lie well ‘‘outside’’ of the tellurium
polyhedron, whereas the copper atoms ‘‘missing’’ the phosphine in 63 are drawn
into the cluster framework. The near linear coordination geometries of the two
two-coordinate Ag centers in 104a are replaced with trigonal planar copper sites in
63, and this is accompanied by an increase in the coordination numbers of two
telluride ligands.
The substituted aryl ligand 2,4,6-Me3 C6 H2 Te (mesTe ) is also suitable for stabilizing polymetallic AgTe frameworks. Thus, in an analogous reaction to that for
104, AgCl:PR3 (PR3 ¼ PEt3 , PnPr3, PEt2 Ph) solutions in ether solvents are reacted
with a combination of Te(mes)SiMe3 and Te(SiMe3 )2 . The clusters [Ag 46 (Temes)12 Te17 (PR3 )16 ] (105) are formed as dark red/black crystals in excellent yield (80–85%)
[71]. Cluster 105a (PR3 ¼ PEt3 ) contains a distribution of telluride ligands (Figure
3-82) in the central portion of the cluster, and tellurolate groups and PEt3 are
on the cluster surface. About the central Te 2 are 12 additional telluride ligands
arranged to form a distorted [non-bonded Te Te ¼ 4.523(2)–5.288(2) Å] centered
icosahedron. The central atom telluride is within the range of bonding distances to
eight silver metal atoms [2.814(2)–3.198(2) Å], which occupy positions within the
Te13 -centered icosahedron, and an additional six Ag atoms reside within its deltahedral faces. This central Ag34 Te13 unit is surrounded by 12 m3 -Te(mes) , 4 Te 2 ,
and 12 Ag centers. The coordination number about the 46 silver atoms in 105a
varies from two (linear) to four (tetrahedral). The more condensed nature of the
framework in 105a is reflected in its color of dark red/black as compared with the
more lightly coloured (red) 104a.
177
178
3 Syntheses and Characterizations
25 °C, 1 eq. PPr3
C7H8
R = C11H23
R’ = Ph
x = 0.5
P∩ P =
Ph2P(CH2)2PPh2
AgO2CR
+
+ R’SeSiMe3
x P∩ P
– SiMe3O2CR
−50 °C, CH2Cl2
R = CH3
R’ = Ph
x=1
P∩ P =
Ph2P(CH2)3PPh2
{[Ag5(TePh)6(Ph2PCH2CH2PPh2)
(Ph2PCH2CH2PPh3)]}∞
106 [72]
[Ag18Te(TePh)15(Ph2P(CH2)3PPh2)3Cl]
107 [72]
[Ag38Te13(TetBu)12(Ph2PCH2CH2PPh2)3]
108 [72]
−50 °C, CH2Cl2
R = CH3
R’ = tBu
x=1
P∩ P =
Ph2P(CH2)2PPh2
The synthesis of tellurido/tellurolato-bridged
silver clusters with bidentate phosphine ligands.
Scheme 3-14
Bidentate phosphine ligands have also been shown to be effective in stabilizing
AgTe polynuclear frameworks, and, as in the case of AgSe complexes, they can
serve as both inter- and intracluster bridges. Unlike the entry into PR3/Ag/Te/TeR
complexes, where AgCl is used as a source of Ag(I), the successful preparation of
{[Ag5 (TePh)6 (Ph2 PCH2 CH2 PPh2 )(Ph2 PCH2 CH2 PPh3 )]}y (106), [Ag18 Te(TePh)15 (Ph2 P(CH2 )3 PPh2 )3 Cl] (107), and [Ag38 Te13 (TetBu)12 (Ph2 PCH2 CH2 PPh2 )3 ] (108)
requires carboxylate salts of Ag(I) (Scheme 3-14) [72].
A suspension of silver laurate and half an equivalent of dppe in toluene can be
solubilized via the addition of one equivalent of tri-n-propylphosphine. The addition of PhTeSiMe3 at room temperature to this solution results in the crystallization of {[Ag5 (TePh)6 (Ph2 PCH2 CH2 PPh2 )(Ph2 PCH2 CH2 PPh3 )]}y (106) in moderate yield (28%) after several days [72]. The structure of 106 (Figure 3-83) consists of
[Ag5 (TePh)6 ] units linked by a dppe ligand. The arrangement of silver centers and
tellurolate ligands is similar to that observed in the copper-tellurium complex
[Cu5 (TePh)6 (PEt3 )3 ][PEt3 Ph]. There are three m3 -TePh groups (AgaTe ¼ 2.786–
2.993 Å) around the ‘‘base’’ of the cluster and three m2 -TePh bridges (AgaTe ¼
2.662–2.837 Å), the Te centers arranged to form a distorted non-bonded Te6 octahedron. There are four silver atoms in the base of the cluster and a lone, trigonal
planar silver atom bonded to the three m2 -TePh bridges. The lone Ag atom centered
in the base of the cluster has noticeably shorter Ag Ag contacts to the other
three basal Ag sites (2.815(7)–2.870(6) Å) than do the other silver atoms in 106
(>3.1 Å). The three additional Ag(I) atoms are tetrahedral, bonded to one PPh2 (CH2 )-group. The dppe ligand serves as an intercluster bridge between two
[Ag5 (TePh)6 ] cluster units, to form the infinite repeat observed in 106. The pen-
3.1 Semiconductor Nanoparticles
Fig. 3-83 A projection of the polymeric repeat in
{[Ag5 (TePh)6 (Ph2 PCH2 CH2 PPh2 )(Ph2 PCH2 CH2 PPh3 )]}y (106).
For clarity, only the ipso carbon atoms of the -PPh3 þ centers
and the C atoms of the methylene units on phosphorus are
illustrated.
107
Fig. 3-84
The core atoms in [Ag18 Te(TePh)15 (Ph2 P(CH2 )3 PPh2 )3 Cl] (107).
dant [Ph2 PCH2 CH2 PPh3 ]þ cation balances the negative charge on the cluster
framework. Although the origin of this cation is difficult to rationalize, the generation of phosphonium centers during photochemically induced condensation reactions of CuaTePh complexes is well established [19, 61].
[Ag18 Te(TePh)15 (Ph2 P(CH2 )3 PPh2 )3 Cl] (107) is synthesized by the 1:1:1 reaction
of AgOAc, Ph2 P(CH2 )3 PPh2 and PhTeSiMe3 in CH2 Cl2 , the highly symmetric
cluster residing on a crystallographic C3 axis (Figure 3-84) [72]. Cluster 107 possesses a pseudo close-packed arrangement of Te centers, with a central telluride
ligand present from TeaCphenyl bond cleavage of a TePh fragment. The five silver
179
180
3 Syntheses and Characterizations
atoms within the range of bonding distances to this central Te 2 are arranged
to form nearly perfect trigonal bipyramidal coordination geometry about the tellurium. The PhTe ligands bridge either three or four Ag atoms, and, consistent
with the rather open nature of the AgaTe framework, the majority (nine) of the 11
interstitial silver atoms are three-coordinate, the other two adopting tetrahedral coordination. All phosphine-bonded Ag(I) sites and the lone AgaCl fragment are also
bonded to three PhTe ligands. The dppp ligands are located exclusively around
the ‘‘base’’ of the pyramidal AgTe framework, with the chloride occupying the apical position on the cluster.
The largest structurally characterized Ag/Te/TeR cluster stabilized with bidentate
phosphine ligands, [Ag38 Te13 (TetBu)12 (Ph2 PCH2 CH2 PPh2 )3 ] (108) is obtained
without the use of Te(SiMe3 )2 as a source of telluride ligands [72]. Cluster 108
forms as small, black crystals in @50% yield from the reaction of AgOAc, dppe and
tBuTeSiMe3 in a 4:4:3 ratio at low temperatures (Scheme 3-14). Gradual warming
of reaction solutions to room temperature leads to the formation of 108 over several weeks, and the stability of these solutions at higher temperature (vs the precipitation of amorphous solids) accounts for the formation of such a ‘‘telluriderich’’ cluster.
The structure of 108 (Figure 3-85) contains a spherical AgaTe framework. The 13
Te 2 ligands are spatially arranged in the cluster core to form a centered icosahedron, as is observed in [Ag 46 (Temes)12 Te17 (PR3 )16 ] (105) [71]. The central telluride
ligand of this deltahedral sub-unit forms eight AgaTe bonding interactions (3.049–
3.199 Å), and the polyhedron is in fact surrounded by twelve Ag(I) atoms, themselves arranged to form an Ag12 (non-bonded) icosahedron. The outer Te12 and inner Ag12 ‘‘shells’’ are then surrounded by the 12 m3 -tBuTe surface ligands and the
additional 26 Ag(I) atoms. The twelve phosphine-bonded silver atoms have tetra-
Fig. 3-85
108
The core atoms in [Ag38 Te13 (TetBu)12 (Ph2 PCH2 CH2 PPh2 )3 ] (108).
3.1 Semiconductor Nanoparticles
hedral coordination as they are each also bonded to three Te centers. Sixteen of the
remaining Ag(I) atoms are trigonal planar, and ten Ag(I) atoms are tetrahedrally
coordinated to four Te ligands in asymmetric fashion.
Selenium-Bridged Gold Clusters
The growing interest in gold-chalcogenide polynuclear chemistry can be attributed
partly to the demonstration that these systems exhibit rich photoluminescence behavior and partly to the interesting structural frameworks that are generated as a
result of weak Au(I) Au(I) interactions [73].
The utility of bidentate phosphine ligands in stabilizing metal-chalcogenide
cores is again demonstrated by the synthesis of [Au10 Se4 (Ph2 PCH2 PPh2 )4 ]Br2
(109) and [Au18 Se8 (Ph2 PCH2 CH2 PPh2 )6 ]Cl2 (110a, Scheme 3-15) from the reaction of Se(SiMe3 )2 , [(AuBr)2 (dppm)] and [(AuCl)2 (dppe)] at ambient temperature,
respectively. The complex dication 109 is illustrated in Figure 3-86 and features
3.1.3.4
[Au10Se4(dppm)4]Br2
109 [74]
[Au18Se8(dppe)6]Cl2
110 [74]
RT, CH2Cl2
X = Br
P∩ P =
Ph2PCH2PPh2
(AuX)2(P∩P)
+ Se(SiMe3)2
– SiMe3Cl
RT, CH2Cl2
X = Cl
P∩ P =
Ph2P(CH2)2PPh2
Scheme 3-15
The synthesis of selenido-bridged gold clusters with bidentate phosphine ligands.
Fig. 3-86 The molecular structure of the cationic framework in
[Au10 Se4 (Ph2 PCH2 PPh2 )4 ]Br2 (109).
181
182
3 Syntheses and Characterizations
four m3-selenide ligands and four bidentate DPPM ligands bonded to ten Au(I)
centers. The m3 -Se 2 ligands are symmetrically bonded to three gold atoms (Sea
Au ¼ 2.423(1)–2.461(1) Å), with two ligands on either side of the Au10 ‘‘plane’’.
The arrangement of the ten gold atoms has eight Au(I) atoms forming an outer,
corrugated octahedron that surrounds two central Au(I) atoms. The molecule can
also be regarded as consisting of two [Au5 Se 2 ]þ units held together by two dppm
ligands together with a contribution from the aurophilic interaction of the central
Au(I) Au(I) atoms (3.127(1) Å). These two gold atoms are bonded to two selenide
ligands, with a noticeable deviation from linear geometry (164.7(1)–168.6(1) ) due
to the metal-metal interaction. The eight additional gold sites are each bonded to
one phosphorus and one selenium ligand, again with moderate deviations from
linear geometries (JPaAuaSe ¼ 168.5(1)–177.2(2) ) and Au(I) Au(I) distances
of 3.031(1)–3.364(1) Å similar to those observed in the structurally related complex
[Au12 (dppm)6 S4 ] 4þ .
The larger gold-selenide cluster [Au18 Se8 (Ph2 PCH2 CH2 PPh2 )6 ]Cl2 (110a) (Figure 3-87) is prepared in 90% yield from [(AuCl)2 (dppe)] and Se(SiMe3 )2 . The
structure of the complex dication 110 consists of two [Au9 (m3 -Se)4 (dppe)3 ]þ fragments formed from three Au 4 Se3 dppe ‘‘rings’’. Each of these forms a catenated
structure with the corresponding ring in the second Au9 fragment such that there
are three pairs of interlocked rings, resulting in a central Au 6 Se 2 heterocubane
unit. The six gold atoms in this heterocubane are two-coordinate, each linearly
bonded to two selenide ligands (JSeaAuaSe ¼ 173.7(1)–174.9(1) ). The Au(I) Au(I) distances in the Au 6 Se 2 unit are in the range 2.940(2)–3.002(2) Å, and in the
rest of the cluster are in the range 2.948(2)–3.327(3) Å. The remaining twelve Au(I)
centers display the same, expected coordination geometries, as they are bonded
to one phosphine and one selenide ligand (JPaAuaSe ¼ 166.8(1)–173.1(1) ). The
m3 -Se ligands symmetrically bridge the Au(I) centers (2.430(2)–2.481(3) Å), with six
A projection of the molecular structure of
[Au18 Se8 (Ph2 PCH2 CH2 PPh2 )6 ] 2þ (110, phenyl rings omitted
for clarity).
Fig. 3-87
3.1 Semiconductor Nanoparticles
of the selenide ligands distributed around the periphery of the AuSe frame. It is
noteworthy that reactions of [(AuCl)2 (dppe)] with RESiMe3 (R ¼ Et, iPr, tBu, nBu)
invariably lead to the formation of [Au18 Se8 (dppe)6 ]Cl2 (110a) via SeaC bond
cleavage, with no evidence to date of any intermediate, AuaSeR, containing complexes that can be isolated. Although the chloride counterions are easily substituted
via the reaction of 110a with KPF6 and NaBPh 4 to yield [Au18 Se8 (dppe)6 ](PF6 )2
(110b) and [Au18 Se8 (dppe)6 ](BPh 4 )2 (110c), respectively.
High-resolution electrospray ionization mass spectrometric investigations of
110b display the expected isotopic distribution centered around m/z 3283.89, corresponding to the dication 110, with no evidence of fragmentation. Controlled,
collision-induced fragmentation reactions of 110 with argon indicate that the
preferred dissociation path is into two AuaSe fragments, [Au13 Se6 (dppe)3 ]þ and
[Au5 Se 2 (dppe)2 ]þ , with the concomitant loss of dppe. These experiments suggest
similar AuaSe and AuaP binding energies in 110.
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(b) M. G. B. Drew, A. H. B. Othman,
D. A. Edwards, R. Richards, Acta
Crystallogr. Sect. 1975, B31, 2695.
44 (a) R. Ahlrichs, M. Bär, M. Häser,
H. Horn, C. Kölmel, Chem. Phys.
Lett. 1995, 242, 652; (b) O. Treutler,
R. Ahlrichs, J. Chem. Phys. 1995,
102, 346; (c) A. Schäfer, H. Horn,
R. Ahlrichs, J. Chem. Phys. 1992, 97,
2571.
45 (a) K. Eichkorn, O. Treutler, H.
Öhm, M. Häser, R. Ahlrichs, Chem.
Phys. Lett. 1995, 242, 652; (b) O. Treutler, R. Ahlrichs, J. Chem. Phys.
1995, 102, 346; (c) K. Eichkorn, F.
Weigend, O. Treutler, R. Ahlrichs,
Theor. Chim. Acta 1997, 97, 119.
46 C. Møller, M. S. Plesset, Phys. Rev.
1934, 46, 618.
47 R. G. Parr, W. Yang, Density
Functional Theory of Atoms and
Molecules, Oxford University Press,
New York 1988.
48 T. Ziegler, Chem. Rev. 1991, 91, 651.
49 E. Lax, Taschenbuch für Chemiker und
Physiker, Vol. I, Springer, Berlin,
Heidelberg 1967, p. 324.
50 E. Lax, Taschenbuch für Chemiker und
Physiker, Vol. I, Springer, Berlin,
Heidelberg 1967, p. 326.
3.2 Synthesis of Metal Nanoparticles
51 A. F. Hollemann, E. Wiberg,
Lehrbuch der Anorganischen Chemie,
W. de Gruyter, Berlin 1985, p. 148.
52 G. Wedler, Lehrbuch der Physikalischen Chemie, VCH, Weinheim 1987,
p. 168.
53 D. Fenske, N. Zhu, T. Langetepe,
Angew. Chem. 1998, 110, 2784; Angew.
Chem. Int. Ed. 1998, 37, 2640.
54 Y. Park, Ph.D. Thesis, Michigan State
University 1992.
55 G. Savelsberg, H. Schäfer, Z.
Naturforsch. 1978, B33, 711.
56 K. O. Klepp, J. Less-Common Met.
1987, 128, 79.
57 G. Savelsberg, H. Schäfer, Mater.
Res. Bull. 1981, 16, 1291.
58 Y. Park, D. C. Degroot, J.
Schindler, C. R. Kannewurf, M. G.
Kanatzidis, Angew. Chem. Int. Ed.
Engl. 1991, 30, 1325.
59 (a) Y. Park, M. G. Kanatzidis,
J. Chem. Mater. 1991, 3, 781;
(b) X. Zhang, Y. Park, T. Hogan,
J. L. Schindler, C. R. Kannewurf,
S. Seong, T. Albright, M. G.
Kanatzidis, J. Am. Chem. Soc. 1995,
117, 10300.
60 M. Emirdag, G. L. Schimek, J. W.
Kolis, J. Chem. Soc., Dalton Trans.
1999, 1531.
61 M. W. DeGroot, M. W. Cockburn,
M. S. Workentin, J. F. Corrigan,
Inorg. Chem. 2001, 40, 4678.
62 (a) F. C. Frank, J. S. Kasper, Acta
Cryst. 1958, 11, 184; (b) F. C. Frank,
J. S. Kasper, Acta Cryst. 1959, 12, 483.
63 (a) C. Kölmel, R. Ahlrichs, J. Phys.
Chem. 1990, 94, 5536; (b) F. A.
Cotton, X. Feng, M. Matusz, R.
Poli, J. Am. Chem. Soc. 1988, 110,
7077; (c) K. M. Merz, R. Hoffmann,
Inorg. Chem. 1988, 27, 2120.
64 U. Müller, Anorganische Strukturchemie, 2 nd Edn., B.G. Teubner,
Stuttgart 1991, p. 101f.
65 (a) Jellinek in D. W. Sharp (Eds.):
MTP International Review of Science,
Inorganic Chemistry Series One, Vol. 5,
Butterworths, London 1972, p. 339;
(b) C. W. Folmer, F. Jellinek, J. Less
Common Met. 1980, 76, 153.
66 R. Ahlrichs, J. Besinger, A. Eichhöfer, D. Fenske, A. Gbureck,
Angew. Chem. 2000, 112, 4089.
67 D. Fenske, N. Zhu, T. Langetepe,
Angew. Chem. Int. Ed. Engl. 1998, 37,
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68 D. Fenske, T. Langetepe, Angew.
Chem. Int. Ed. Engl. 2002, 41, 300.
69 J. F. Corrigan and D. Fenske, Chem.
Commun. 1996, 943.
70 J. F. Corrigan, D. Fenske, W. P.
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71 J. F. Corrigan, D. Fenske, Chem.
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72 T. Langetepe, D. Fenske, Z. Anorg.
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73 (a) H. Schmidbaur, Gold: Progress in
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173 and references therein.
74 D. Fenske, T. Langetepe, M. M.
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Chem. Int. Ed. Engl. 2000, 39, 1857.
3.2
Synthesis of Metal Nanoparticles
Section 3.2 is divided into two parts, the first of which considers particles mainly
consisting of noble metals, except for a few examples. Nanoparticles of noble metals have an especially long tradition, and fundamental synthetic strategies have
therefore been known for some decades. The purpose of this part is therefore not
to describe the huge number of synthetic pathways to metal nanoparticles, but
rather to summarize the known syntheses. However, novel procedures are included if they complement traditional ones.
185
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3 Syntheses and Characterizations
The second part of Section 3.2 deals exclusively with magnetic particles, since
these species have attracted a lot of attention during the last decade. The reason is
to be seen in the manifold future applications in nanodevices. In contrast to the
part dealing with noble metal nanoparticles, this second part will also describe the
most important properties of the particles.
3.2.1
Noble Metal Nanoparticles
John S. Bradley and Günter Schmid
Introduction
The advancement of the understanding of small-particle science and the potential
for new materials science based on the chemistry and physics of nanoscale metal
clusters rests on the measurement and application of useful size-dependent properties of small metal nanoparticles. Ideally this requires the preparation and isolation of monodisperse metal particles with a great degree of control over size,
structure and composition. Metals in varying levels of dispersion have a long history in many technological applications, but the methods of preparation used are
usually determined by the intended use for the metal and are not necessarily developed with the aim of preparing well-defined metal particles. For example, highly
dispersed metal catalysts are usually prepared by high-temperature reduction of
metal precursors on oxide supports. These severe preparative conditions tend to
yield metal particles with ranges of sizes and morphologies broad enough to hide
the potentially unique properties which a monodisperse nanoscale sample could
exhibit.
With these considerations in mind, synthetic chemists have begun to address the
needs of metal particle research by developing the synthetic chemistry of colloidal
metals, with a view to using the strategies of molecular chemistry to prepare welldefined metal particles. The goal may be stated as the search for synthetic methods
for preparing metal particles stabilized in the 1–20 nm size range, and these synthesis approaches will be covered in this chapter. This area was discussed in a
recent book giving a broad review of metal colloid chemistry [1], and is brought up
to date here with a particular focus on synthetic methods.
3.2.1.1
History and Background
Metal colloid science (as opposed to alchemical and aesthetic investigations) can be
said to have begun with the experiments of Michael Faraday on gold sols in the
mid-nineteenth century [2]. Deep red solutions of colloidal gold were prepared by
the reduction of chloroaurate [AuCl 4 ] solutions using phosphorus as a reducing
agent. (Recently these preparations have been reproduced in Faraday’s original
laboratory at the Royal Institution in London by J. M. Thomas, who has demonstrated that the gold sols thus produced contain particles of diameter 3–30 nm [3].
The work of Zsigmondy at the turn of the century put the physical investigation of
3.2.1.2
3.2 Synthesis of Metal Nanoparticles
colloidal solutions (or suspensions – he used the terms interchangeably but with
frequent comment as to the most appropriate term) on a firm basis, and his invention of the ultramicroscope enabled careful studies of the effect of preparation
conditions on particle size to be carried out [4]. Many methods for the preparation
of colloidal metals have been developed over the years since Faraday’s experiments.
A summary of the research on the preparation of metal sols in the first half of the
century can be found in the 1951 paper of Turkevitch and coworkers [5], which
describes detailed synthesis conditions and early electron microscopic characterization of gold sols prepared by the methods of Faraday (see above), Bredig (electric
arc between gold electrodes under water), and Donau (action of carbon monoxide
on chloroauric acid), and by a number of other related methods involving the
action of reducing agents on chloroaurate solutions. These preparations introduce
many of the methods used today in preparative metal colloid chemistry. Turkevich’s
long study of this system yielded detailed knowledge of the nucleation, growth, and
agglomeration of colloidal metals and has also allowed the reproducible preparation of gold sols with narrow particle size distributions in a range of mean sizes
[6]. Figure 3-88 shows a TEM image of typical, very narrowly size-distributed 18
nm gold nanoparticles prepared by the Turkevitch method.
In general, the particle sizes found for these and related colloids can be varied
over the range ca. 2–100 nm, depending on preparative conditions, with most examples falling in the large diameter range.
Inorganic colloid chemistry in the last two decades has focused on inorganic
chalcogenides [7, 8] in approaches to quantum-confined semiconductor particles,
and has progressed to the point where truly monodisperse II–VI semiconductor colloids in size ranges from 12 to 115 Å have been isolated and their sizedependent optical and other physical properties measured [9, 10]. The synthetic
chemistry of metal colloids has now reached this degree of refinement, and after a
lengthy period in which highly dispersed metals were widely exploited in specific
technological applications such as photography and catalysis, the resurgence of
interest in the intrinsic chemistry and physics of small metal particles has fostered
the establishment of a new synthetic metal colloid chemistry.
Generation of Metal Particles
The successful establishment of this new area of inorganic materials chemistry
will depend first of all on the control of particle size, morphology and composition
of metal nanoparticles in syntheses which produce chemically significant quantities. So, for our purposes, the preparation of size-selected particles in a molecular
beam experiment, although elegant, is not relevant to materials chemistry. The
problem of preparing isolable yields of uniform samples is also one which was
central to the development of synthetic organometallic cluster chemistry. The obvious difference between molecular synthesis and molecular approaches to colloidal metals is that with the former it is the constant aim of the organometallic
chemist to maintain metals in complexed molecular form, whereas with the latter
the goal is reversed, and we seek controlled methods for metal formation from
molecular or ionic precursors. A reaction which results in the formation of ele3.2.1.3
187
188
3 Syntheses and Characterizations
Fig. 3-88 Transmission electron micrograph of a
monodisperse 18 nm gold hydrosol prepared by citrate
reduction of HAuCl4 .
mental metal in the laboratory of the synthetic molecular chemist is a failure, but it
is also a potential source of colloidal metal. Such a process would normally lead to
the formation of a metallic precipitate or mirror, but, in the presence in solution of
an appropriate stabilizing agent, aggregation of the metal atoms can be arrested at
an early stage, and the metal particles are preserved in suspension in the organic
liquid.
A survey of the literature reveals that many of the various methods reported for
the preparation of metal nanoparticles are applicable to a number of metals across
the periodic table. For example, salt reduction using main group hydride reducing
agents has been used for the preparation of many metals in colloidal form, not
only those cited in this review. It is not the goal of this chapter to provide a direc-
3.2 Synthesis of Metal Nanoparticles
tory of all reports of colloid syntheses, but only to give examples of the principal
types of preparative methods which can be used.
Stabilization of Colloidal Metal Particles in Liquids
Before beginning a description of synthetic methods, a general and crucial aspect
of colloid chemistry should be considered, and that is the means by which the
metal particles are stabilized in the dispersing medium, since small metal particles
are unstable with respect to agglomeration to the bulk. At short interparticle distances, two particles would be attracted to each other by van der Waals forces and
in the absence of repulsive forces to counteract this attraction an unprotected sol
would coagulate. This counteraction can be achieved by two methods, electrostatic
stabilization and steric stabilization [11]. In classical gold sols, for example, prepared
by the reduction of aqueous [AuCl 4 ] by sodium citrate, the colloidal gold particles
are surrounded by an electrical double layer formed by adsorbed citrate and chloride ions and cations which are attracted to them. This results in a Coulombic repulsion between particles, and the net result is shown schematically in Figure 3-89.
The weak minimum in potential energy at moderate interparticle distance defines a stable arrangement of colloidal particles which is easily disrupted by medium effects and, at normal temperatures, by the thermal motion of the particles.
Thus, if the electric potential associated with the double layer is sufficiently high,
electrostatic repulsion will prevent particle agglomeration, but an electrostatically
3.2.1.4
Fig. 3-89 Electrostatic stabilization of metal colloid particles.
Attractive van der Waals forces are outweighed by repulsive
electrostatic forces between adsorbed ions and associated
counterions at moderate interparticle separation.
189
190
3 Syntheses and Characterizations
Fig. 3-90 Aggregation of an electrostatically stabilized 12–
18 nm gold hydrosol caused by the addition of pyridine, a
neutral ligand (reproduced with permission from [12]).
stabilized sol can be coagulated if the ionic strength of the dispersing medium is
increased sufficiently. If the surface charge is reduced by the displacement of adsorbed anions by a more strongly binding neutral adsorbate, the colloidal particles
can now collide and agglomerate under the influence of the van der Waals attractive forces. This phenomenon is nicely demonstrated by the addition of pyridine to
a gold sol of the type mentioned above, as shown in Figure 3-90 [12].
Even in organic media, in which electrostatic effects might not normally be considered to be important, the development of charge has been demonstrated on inorganic surfaces, including metals, in contact with organic phases such as solvents
and polymers [13]. For example, the acquisition of charge by gold particles in organic liquids has been demonstrated, and the sign and magnitude of the charge
has been found to vary as a function of the donor properties of the liquid [14].
Thus, even for colloidal metals in suspension in relatively non-polar liquids, the
3.2 Synthesis of Metal Nanoparticles
possibility cannot be excluded that electrostatic stabilization contributes to the stability of the sol.
A second means by which colloidal particles can be prevented from aggregating
is by the adsorption of molecules such as polymers, surfactants or ligands at the
surface of the particles, thus providing a protective layer. Polymers are widely used,
and it is obvious that the protectant, in order to function effectively, must not only
coordinate to the particle surface, but must also be adequately solvated by the dispersing fluid – such polymers are termed amphiphilic. The choice of polymer is
determined by consideration of the solubility of the metal colloid precursor, the
solvent of choice, and the ability of the polymer to stabilize the reduced metal particles in the colloidal state. Natural polymers such as gelatin and agar were often
used before the advent of synthetic polymer chemistry, and related stabilizers such
as cellulose acetate, cellulose nitrate [15], and cyclodextrins [16] have been used
more recently. Thiele [17] proposed the Protective Value as a measure of the ability
of a polymer to stabilize colloidal metal. It was defined, similarly to the older Gold
Number of Zsigmondy, as the weight of the polymer which would stabilize 1 g of a
standard red gold sol containing 50 mg/L gold against the coagulating effect of 1%
sodium chloride solution. Several other studies have been performed on the relative ability of polymers to act as steric stabilizers [18–20], and, despite the fact that
these quite subjective studies focus on very specific (and quite different) sol systems, it seems that, of the synthetic polymers considered, vinyl polymers with
polar side groups such as poly(vinylpyrrolidone) (PVP) and poly(vinyl alcohol) are
especially useful in this respect.
The use of copolymers introduces another degree of variability to colloidal stabilization, as the comonomer ratio can be varied. For example, the use of vinylpyrrolidone-vinyl alcohol copolymers is reported for the preparation of platinum and
silver hydrosols [21]. The silver sols were stable only in the presence of the copolymer, and the size of the silver particles decreased with an increase in vinylpyrrolidone content of the copolymer.
Electrostatic and steric stabilization are in a sense combined in the use of longchain alkylammonium cations and surfactants, either in single-phase sols or in
reverse micelle synthesis of colloidal metals. Examples will be described below. A
new class of metal colloids has recently been established in which the surface of
the particle is covered by relatively small ligand molecules such as sulfonated triphenylphosphine [22] or alkane thiols, and examples will be given below [23].
Synthetic Methods for the Preparation of Colloidal Transition Metals
The synthetic methods which have been used include modern versions of established methods of metal colloid preparation such as the mild chemical reduction of
solutions of transition metal salts and complexes and newer methods such as
radiolytic and photochemical reduction, metal atom extrusion from labile organometallics, and the use of metal vapor synthesis techniques. Some of these reactions
have been in use for many years, and some are the results of research stimulated
by the current resurgence in metal colloid chemistry. The list of preparative methods is being extended daily, and, as examples of these methods are described be3.2.1.5
191
192
3 Syntheses and Characterizations
low, the reader will quickly be made aware that almost any organometallic reaction
or physical process which results in the deposition of a metal is in fact a resource
for the metal colloid chemist. The acquisition of new methods requires only the
opportunism of the synthetic chemist in turning a previously negative result into a
synthetic possibility.
Salt Reduction
The reduction of transition metal salts in solution is the most widely practiced
method of generating colloidal suspensions of the metals. In aqueous systems, the
reducing agent must be added or generated in situ, but in non-aqueous systems the
solvent and the reducing agent can be one and the same. Easily oxidized solvents
such as alcohols can thus function both as reducing agents and as the colloid diluent, and these have been used widely in colloid preparations. Hirai and Toshima
developed the so-called ‘‘alcohol reduction process’’, which is a very general process
for the production of metal nanoparticles, often stabilized by organic polymers
such as poly(vinylpyrrolidone), poly(vinyl alcohol), and poly(methyl vinyl ether)
[24–31]. In general, the alcohols which were useful reducing agents contained ahydrogen (tert-butanol, for example, was not effective) and were oxidized to the
corresponding carbonyl compounds (methanol to formaldehyde, ethanol to acetaldehyde, iso-propanol to acetone) which were detected during the reaction [Eq. (1)].
RhCl3 þ 3/2 R1 R2 CHOH ! Rh(0) þ 3/2 R1 R2 CbO þ 3 HCl
ð1Þ
Palladium acetate has been compared with palladium chloride as a starting material for PVP-stabilized colloidal palladium and was found to give smaller particles
and better control of particle size distribution, and the effect of Pd/PVP ratio was
also investigated [32]. Palladium organosols have been prepared by the reduction
of palladium acetate in refluxing methanol or at a higher temperature in 2-ethoxyethanol, both in the presence of PVP, which yields ca. 6.5 nm particles [33, 34].
Organosols of platinum and palladium have been synthesized by heating salts of
the metals in long-chain aliphatic alcohols like 1-decanol, which not only reduce
the salts to the metal but also act as stabilizers for the colloid particles [35]. Esumi
has reported the use of several alcohols in a comparative synthetic study in which
C1 aC5 aliphatic alcohols, a-phenylethanol, b-phenylethanol, and benzyl alcohol
were used as solvent and reducing agent for palladium acetate in the presence of
PVP, giving particle sizes in the 1–4 nm range [36]. A review reports on the generation of bimetallic nanoparticles [24].
The reduction of metal salts by added reducing agents in non-reducing solvents
is the oldest established procedure for the preparation of aqueous suspensions of
colloidal metals. A wide range of reducing agents have been used in metal colloid
syntheses, and they are frequently interchangeable from metal to metal. Thus, hydrazine hydrate has been used in preparations of colloidal platinum [37], gold [38],
copper [39], and palladium-platinum alloy [40]. A catalog of preparations will
not be presented here, but examples of the use of typical reducing agents will be
given.
3.2 Synthesis of Metal Nanoparticles
Faraday reduced aqueous solutions of [AuCl 4 ] with phosphorus vapor to produce gold hydrosols [2], and a variety of reducing agents have subsequently been
used with tetrachloroaurate, both with and without protective polymers, to give
colloidal gold with particle sizes from one to several hundred nanometers in diameter. Turkevitch and his coworkers reproduced many of these methods [5] and
have established reliable preparative procedures for the preparation of gold sols
with quite precisely defined particle sizes. The 20 nm gold sol prepared by the
reduction of [AuCl 4 ] by sodium citrate has become a standard for histological
staining applications. The mechanism of the reduction of metal salts by citrate was
also studied by the same authors [5], who demonstrated that, in the reduction of
[AuCl 4 ] , an induction period which was present when citrate was the reducing
agent was absent when [AuCl 4 ] was reduced by acetone dicarboxylate (an oxidation product of citrate), and a rapid formation of colloidal gold resulted.
CH2 -COOH
HOC-COOH
CH2 -COOH
CH2 -COOH
½AuCl4 ƒƒƒ
ƒ! ObC
ƒƒƒƒ! CO2 þ Au(0)
ð2Þ
CH2 -COOH
A similar mechanism may be postulated for other reductions with citrate, and this
method has been widely used for the preparation of other metal sols, such as platinum [41–43], and can be added as an ionic stabilizer in preparations which require an additional reducing agent. The use of formate, citrate, and acetone dicarboxylate as reducing agents at various pH values was reported to give good control
over particle size in the preparation of a series of platinum hydrosols [44].
A study by Glaunsinger and coworkers [37] has been reported on the reduction
of chloroplatinic acid by a variety of reducing agents including sodium borohydride, hydroxylamine hydrochloride, dimethylamine borane, sodium citrate, hydrazine monohydrate, sodium formiate, trimethylamine borane, sodium trimethoxyborohydride, and formaldehyde.
Among the various reducing agents, BH4 plays a dominant role. Brust et al.
[45] introduced this method to reduce gold salts in the presence of alkanethiols,
resulting in 1–3 nm Au particles. Variations in the thiol concentrations provide
good size control of the particles, between 2 and 5 nm [46]. Thiol-stabilized Au
particles became available by following a seeding growth approach starting with 3.5
nm colloids [47]. Nanoparticles of gold in the size range of 10 to about 30 nm have
become known using mercaptosuccinic acid for stabilization and BH4 as the reducing species. Variations in the ratio [AuCl 4 ] : succinic acid allow the specific
synthesis of 10.2, 10.8, 12.8, 19.4, and 33.6 nm colloids [48].
The generation of a reducing agent in situ was recently used for the preparation of gold hydrosols from [AuCl 4 ] . Tetrakis(hydroxymethyl)phosphonium chloride (THPC) hydrolyzes under alkaline conditions to eliminate formaldehyde, and
193
194
3 Syntheses and Characterizations
under hydrolysis conditions [AuCl 4 ] oxidizes THPC with the formation of small
(1–2 nm) colloidal gold clusters [49–51]. The mechanism of this reaction is not
known, but it is known that under high pH conditions THMP yields formaldehyde
and subsequently hydrogen, either of which could be the reducing agent.
Hydrogen has been used effectively as a reducing agent for metal salts in electrostatically stabilized metal sol syntheses. Hydrosols of palladium, platinum, rhodium, and iridium, stabilized with poly(vinyl alcohol), were prepared by hydrogen
reduction of the metal hydroxides [52–56].
Moiseev et al. used hydrogen to synthesize Pd nanoclusters of very uniform size
and composition, so that these particles can be described by a rather precise formula: Pd561 LA60 (OAc)A180 (L ¼ phenanthroline, bipyridine) [57–62]. Pd nanoclusters, also of discrete size, shape, and composition have been described by
Schmid et al., also using H2 to reduce Pd salts: Pd561 phen36 OA200 , Pd1415 phen60 OA1100 , Pd2057 phen84 OA1600 (phen ¼ phenanthroline) [63–65].
These nanoclusters belong to the family of so-called full-shell clusters, meaning
that their metal nuclei consist of a number of atoms that follows hexagonal or
cubic close packing. The coordination number 12 in such metallic structures leads
to a first full-shell cluster of 1 þ 12 ¼ 13 atoms, followed by another shell of 42
atoms, resulting in 55 atoms, etc. Generally, a nanocluster of n shells consists of
10n 2 þ 2 atoms. So, the above-mentioned Pd clusters correspond to five shells
(561), seven shells (1415), and eight shells (2057). Of course, with increasing size,
the exact number of atoms can in practice deviate from the ideal number. However, it has to be stated that the full-shell clusters that have become known up to
the present time are indeed more or less monodisperse, especially compared to the
huge number of larger nanoparticles, which usually show size distributions of at
least G10%.
The Au13 and Au55 clusters at the very beginning of this row of full-shell clusters
have to be discussed. Au13 clusters were first described in 1981 [66, 67] and were
found to form icosahedral structures, whereas the two-shell clusters with Au55
cores exhibit cubic close-packed (ccp) structures like bulk gold.
Fig. 3-91
Model of Au55 (PPh3 )12 Cl6 .
3.2 Synthesis of Metal Nanoparticles
Au55 (PPh3 )12 Cl6 is probably the most frequently investigated full-shell cluster,
since its size definitely represents the transition from bulk to molecule (represented by ligand-protected Au13 clusters), so that it represents the prototype metallic quantum dot [68–117]. Chapters 4 and 5 deal intensively with this cluster and
some of its ligand-modified relatives. The synthesis of Au55 (PPh3 )12 Cl6 occurs by a
very special reduction method. B2 H6 is used to react with Ph3 PAuCl in warm
benzene or toluene [68, 69]. Diborane not only acts as a reducing agent but also
traps excess PPh3 to give Ph3 P-BH3 . Numerous derivatives of the PPh3 -stabilized
cluster are available by ligand exchange reactions. For instance, Ph2 PC6 H4 SO3 Na
removes PPh3 quantitatively and makes the compound water-soluble [73]. Thiols
are best suited to substitute PPh3 because AuaS bonds are stronger than AuaP
bonds [102, 103, 118]. The only four-shell cluster that has become known is
Pt309 phen 36 O30G10 (phen ¼ 4,7-C6 H4 SO3 Na-substituted 1,10-phenanthroline)
[119]. It is prepared from Pt(II) acetate, which is reduced by hydrogen at room
temperature in acetic acid solution in the presence of equivalent amounts of ligand. Under these conditions, the oxygen-free cluster is formed first, and this is
then carefully treated with oxygen to give the air-stable final product. The nature of
the oxygen is probably (as in the above-mentioned Pd full-shell clusters) superoxidic or peroxidic. It can easily be removed from the surfaces of the clusters by
hydrogen. The structure of the Pt309 core is ccp, as in the case of the Au55 , the
Pd561 , the Pd1415 , and the Pd2057 nanoclusters, as has been shown by X-ray diffraction and high-resolution electron microscopic (TEM) studies. More structural details are described in [120].
A remarkable type of rather well-defined nanoclusters has been described by
Finke et al. [121–123]. Polyoxoanion-supported Ir(I) complexes, e.g., (Bu 4 N)5 Na3 [(1,5-COD)IrP2 W15 Nb3 O62 ], are reduced in solution by hydrogen to give Ir nanoclusters with Ir cores containing ca. 300–900 atoms, protected by an assembly of
polyoxoanions.
Carbon monoxide/water was used to reduce PtO2 to nanoparticles, effectively
using a water gas shift reaction to provide reducing equivalents of electrons from
the reaction [124].
CO þ H2 O ! CO2 þ 2Hþ þ 2e
ð3Þ
Remarkable contributions to the synthetic field of transition metal nanoparticles
came from Bönnemann’s group. Their work has recently been reviewed [31]. Some
details will be discussed in the following. Tetraalkylammonium hydrotriorganoborates [R 4 N]þ [BR 0 3 H] as reducing species for transition metal salts turned out
to result in numerous valuable [NR 4 ]þ -stabilized metal nanoparticles [31, 125–
128]. Equation (4) represents the general formation of metal nanoparticles following this route.
MXn þ n[R 4 N][BEt3 H] ! M(0) þ n[R 4 N]X þ nBEt3 þ n/2 H2
(X ¼ Cl, Br; R ¼ alkyl, C6 aC20 ; M ¼ metals of Groups 6–11)
ð4Þ
195
196
3 Syntheses and Characterizations
The ammonium salts accomplish two functions: the hydride reduces the positively
charged metal in MXn and the [R 4 N]þ cations stabilize the metal nanoparticles.
The [BEt3 H] anion as the potassium salt even reduces salts of the early transition
metals, only stabilized by THF. So, Ti, V, and Nb colloids could be obtained [123].
A rather unexpected reaction happens between metal salts MXn (M ¼ metals of
Groups 6–11, X ¼ halogen, acetylacatonate, n ¼ 2–4, R ¼ C1 aC8 alkyl) and AlR3 ,
giving particles in the size range 1–12 nm [129, 130]. It is suggested that a not
precisely determinable layer of organoaluminum species covers the nanoparticles.
Silanes have also found use in colloid preparation, and an interesting example of
the use of organosilanes in the formation of organosols of transition metals is the
reduction of platinum, palladium, and rhodium salts to colloidal particles by trialkoxysilanes and trialkylsilanes. This was reported by Lewis and coworkers [131–
133], who demonstrated that presumably homogeneous hydrosilylation catalysts
formed from platinum compounds in the presence of silanes were in fact colloidal
in nature, raising the prospect of colloid activity in many ostensibly homogeneous
catalyst systems. Silanes and SiaH-containing polymers have been used to reduce
Pd(hfacac)2 (hfacac ¼ hexafluoroacetylacetonate) to colloidal palladium in organic
solvents [134].
Alkanethiols have been used as phase transfer stabilizers for gold colloids. The
thiols were added to THF-water solutions of a preformed gold sol (prepared by the
citrate method) from which adsorbed anions had been removed by treatment with
an ion exchange resin, resulting in a cyclohexane-soluble sol [23]. This was a particularly useful procedure in that the colloidal gold particles of various sizes could
be deposited electrophoretically in ordered close-packed two-dimensional lattices
with spacings correlating to the length of the alkyl chain of the thiol.
Surfactants such as water soluble-phosphines, e.g., sodium salts of sulfonated
triphenylphosphine, P(m-C6 H4 SO3 Na)3 , and alkyltriphenyl-methyltrisulfonates,
RC( p-C6 H4 SO3 Na)3 , function as stabilizers for hydrosols of colloidal rhodium,
prepared by hydrogen reduction of aqueous RhCl3 3H2 O in the presence of the
surfactant [135, 136]. The colloidal material is described as a polyhydroxylated
rhodium particle, implying considerable oxidation of the rhodium surface to Rh(I),
and is stabilized by interaction between the anionic head of the surfactant and the
Rh(I) sites at the colloid surface. Conversely, if the rhodium sol is prepared by
aqueous borohydride reduction of RhCl3 3H2 O, the colloidal particles are thought
to be zero-valent rhodium metal, stabilized by interaction with the hydrophobic
alkyl group on the surfactant [137]. Analogous colloids of gold were prepared by
Schmid by citrate reduction of chloroaurate followed by addition of (C6 H5 )2 P- pC6 H5 SO3 Na, giving a water-soluble gold colloid which can be isolated as a solid
and redissolved without aggregation [22].
Just as the use of a surfactant salt of triethylborohydride introduces reducing
agent and stabilizer in a single reagent [125–128], polymeric reducing agents can
be used which also can function as polymeric stabilizers. Examples include the use
of polyethyleneimine in the preparation of colloidal silver [17, 138].
A chemical analog of the preparation of metal colloids by the reduction of metal
salts by radiolytically produced radicals and electrons (see below) can be found in
3.2 Synthesis of Metal Nanoparticles
recent reports by Dye et al. [139, 140], who used alkalides and electrides as homogeneous reducing agents to reduce main group and transition metal salts in
aprotic solvents. These reducing agents are crystalline salts of alkali metal anions
or trapped electrons of the form Kþ [(crown)2 K] or Kþ [(crown)2 e] prepared by
reaction of the alkali metal with a crown ether. In solvents such as dimethyl ether
or THF these remarkable compounds reduce metal salts to the colloidal metal
state with particle sizes in the nanoscale range. For example, AuCl3 is reduced to
an unstable colloidal gold by Kþ [(15-crown-5)K] in dimethyl ether. The familiar
stable ruby red hydrosol is produced by dissolving the reaction product in water.
Electrons have been used in the more traditional form in the electrolytic reduction of metal salts to metal nanoparticles [141–147]. The relevant metal is dissolved
from the anode, and reduction at the cathode in the presence of tetraalkylammonium salts results in the formation of metal nanoparticles. The most significant
aspect of this method is that the size of the resulting metal particles is determined
by the current density, with the smallest particles (1.4 nm) being produced at the
highest current density.
Ti, Fe, Co, Ni, Pd, Pt, Ag, and Au nanoparticles became available in preparative
amounts by this method with yields of >95%. The electrochemical method also
opens pathways to bimetallic nanoparticles such as Pd/Ni, Fe/Co, and Fe/Ni by
using two sacrificial Mbulk anodes [145].
Thermolysis
Since many organometallic compounds of transition metals decompose thermally
to their respective metals under relatively mild conditions, these compounds provide a rich source of nanoparticle precursors. The method is widely applicable. The
thermolysis of carbonyl-containing complexes of rhodium, iridium, ruthenium,
osmium, palladium, and platinum in polymer solutions has been used to prepare
polymer-stabilized colloidal metals with particle sizes in the range 1–10 nm, possibly by decomposition of polymer-bound organometallic intermediates [148–152].
Organosols of palladium and platinum [153–158] have been synthesized by the
thermolysis of precursors such as palladium acetate, palladium acetylacetonate and
platinum acetyacetonate in high boiling organic solvents such as methylisobutylketone, and bimetallic colloids of copper and palladium have been prepared by the
thermolysis of mixtures of their acetates in similar solvents. These preparations
were performed in the absence of stabilizing polymers, and as a result relatively
broad size distributions and large particles were observed. A preparation of bimetallic PVP-stabilized copper-palladium particles with diameters of 1–4 nm has
been reported by Toshima, in which palladium acetate and cupric sulfate are
transformed first to the colloidal hydroxides by treatment of an ethylene glycol
solution of these salts with sodium hydroxide solution in the presence of PVP
[159]. The colloidal mixed hydroxide was then heated to 198 C under nitrogen,
yielding a stable bimetallic sol.
In view of the extensive literature on the syntheses and structures of bimetallic
metal carbonyl clusters, it is even more surprising that few attempts have been
reported to use the facile decomposition of these compounds in the preparation of
197
198
3 Syntheses and Characterizations
bimetallic colloids. Although this would be an unnecessarily complicated route for
colloids of miscible metals if their salts could be coreduced to the same end, there
remains the possibility of preparing bimetallic particles of immiscible metals
starting from well-defined molecular bimetallic clusters of immiscible metals.
Such clusters are known to exist despite the immiscibility of their constituent
metals in the bulk, and there is reason to believe that bulk immiscibility might not
apply to very small particles because of the importance of surface effects.
Photolysis and Radiolysis
Photochemical colloid syntheses fall into two categories: the reduction of metal
salts by radiolytically produced reducing agents such as solvated electrons and free
radicals, and the photolysis of photolabile metal complexes. Again, the essence of
the preparative procedure is the generation of zero-valent metal under conditions
which prevent or at least retard the formation of bulk metal precipitates.
The method has of course a long and important history in the formation of
photographic images in silver halide emulsions. Research over the past decades,
predominantly that of Henglein and of Belloni, has established a wide range of
colloidal metals prepared by this method, encompassing both main group metals
such as cadmium, thallium, and lead, and other noble and non-noble transition
metals [160–175]. Radiolytic methods differ in the identity of the reducing species
formed under irradiation, this being a function of solvent and any added solute.
The radiolysis of aqueous solutions of metal ions produces solvated electrons,
which may react either with the dissolved metal ions directly or with other solutes
to produce secondary radicals, which then reduce the metal cation. For example,
irradiation of solutions of Cu(ClO4 )2 containing HCOONa yields colloidal copper
[173] via reduction of Cu 2þ by both solvated electrons and CO2 :
eðsolvÞ;CO2 Cuþ
nCuþ
Cu 2þ ƒƒƒƒƒƒ! Cu(CO)2 ƒƒƒƒƒƒ! Cu2 þ ƒƒƒƒƒƒ! nuclei ƒƒƒƒƒƒ! Cu(0)
ð5Þ
An advantage of radiolytic methods of colloidal metal synthesis lies in the fact that
a large number of metal nuclei are produced homogeneously and instantaneously,
a condition favorable for the formation of very highly dispersed particles. In addition, the reducing agents are formed in situ, and thus chemical compatibility of
added reducing agents (BH4 , Et3 BH , etc.) with other components of the colloid
system is not an issue, although the possibility of photolysis of those components
exists, as does the possibility of their reaction with primary photolysis or radiolysis
products.
The photolysis of metal complexes and salts is also an effective and clean
method for the generation of colloidal metals, and it is amenable to all of the
experimental protocols used for stabilization of the resulting particles. For example, gold, silver or platinum salts may be subjected to UV-visible irradiation in inverse microemulsions, surfactant solutions or polymer solutions [176–181]. In a
comparative study of hydrogen reduction vs photoreduction of chlorplatinic acid in
the presence of surfactants it was found that photoreduction produced smaller
platinum particles with a narrower size distribution [176].
3.2 Synthesis of Metal Nanoparticles
Ru
3 bar H2
solvent
45 minutes
room temp.
Ru
+
2
[Ru]n
Scheme 3-16
Formation of Ru nanoparticles from Ru(COD)(COT).
Ligand Reduction and Ligand Displacement from Organometallics
Reduction of the metal can be carried out prior to colloid preparation, giving a
zero-valent metal complex as the immediate colloid precursor. The synthesis of
metal carbonyls and their subsequent thermolysis in nanoparticle synthesis is an
example of this approach. In a further refinement, the supporting ligands in the
zero-valent complex can be removed by ligand displacement with an excess of a
weakly bound ligand or by ligand reduction, thus generating metal particles at low
temperatures. The organometallic literature is replete with examples of reactions of
organometallic complexes with hydrogen to give metal deposits, and some of these
reactions have been applied to colloid synthesis.
The zero-valent palladium and platinum complexes with dibenzylideneacetone
Pd(dba)2 and M2 (dba)3 (M ¼ Pd, Pt) have been known since 1970 to react under
mild conditions with either hydrogen or carbon monoxide with the formation of
metal [182].
Meanwhile, there exists a long series of examples where CO and H2 are used to
decompose organometallic precursor molecules [183–200]. As a very recent example, the decomposition of Ru(COD)(COT) (COD ¼ cyclooctadiene, COT ¼ cyclooctatriene) in an atmosphere of hydrogen should be mentioned [201, 202].
The precursor molecule is dissolved in a methanol/THF mixture and is contacted with H2 at 3 bar and room temperature for at least 45 min. Depending on
the MeOH/THF mixture, the Ru particle size can be set at a value between 3 and
86 nm.
3.2.2
Synthesis and Characterization of Magnetic Nanoparticles
Dmitri V. Talapin, Elena V. Shevchenko, Horst Weller
Introduction
Magnetic particles of micro- and nanometer size are attracting growing fundamental and technological interest. Magnets are important for motors, electrical
power transformation, data storage devices, electronics, etc. Miniaturization, one of
the important features of modern technology, requires extremely small magnets of
3.2.2.1
199
200
3 Syntheses and Characterizations
micro- or even nanometer size. How to prepare such materials? Which properties
can one expect from an extremely small piece of a material which is ferromagnetic
in the bulk state? In this Section we will focus on the nanoparticles of ferromagnetic metals, their preparation, and their structural and magnetic characterization.
In nature there are only three ferromagnetic elements at room temperature: iron,
cobalt and nickel (one can also include in this series gadolinium which is ferromagnetic below 16 C). Additionally, some compounds and alloys of manganese,
chromium, and europium can exhibit ferro- or ferrimagnetic behavior [203]. The
exciting discovery that rhodium clusters (Rhn with n ¼ 12–32) are ferromagnetic
[204] shows the non-trivial scaling behavior of magnetism and demonstrates that
magnetic nanomaterials can reveal novel unexpected phenomena.
Magnetic metal nanoparticles could find future applications also as catalysts [205]
and single-electron devices [206–209]. Monodisperse magnetic nanoparticles can
organize themselves into highly ordered superstructures and even form macroscopic ‘‘crystals’’ with nanoparticles as building blocks. The electron spins inside a
nanoparticle can be aligned in a certain direction by applying an external magnetic field. Thin granular films of ferromagnetic nanocrystals are already the basis
of conventional magnetic storage media (hard drives). It is expected that the advanced magnetic media based on ordered arrays of monodisperse nanocrystals
could reach magnetic recording densities in the range of 100 Gb/in 2 and 1 Tb/in 2
[210]. Spin-dependent tunneling electron transport has recently been demonstrated in an array of close-packed cobalt nanoparticles [206]. Granular materials
consisting of nanometer-sized magnetic particles in a dielectric or nonmagnetic
metal matrix exhibit outstanding magnetoresistive properties [211–215].
Magnetic colloids, or ferrofluids, are based on magnetic nanoparticles coated
with a layer of surfactant dispersed in a carrying solvent. The molecules of surfactant (further referred to as ‘‘stabilizing agents’’ or ‘‘stabilizers’’) bind to the nanoparticle surface preventing coagulation and providing solubility and desired surface properties. Moreover, stabilizers protect the particle from oxidation and can
strongly affect the magnetic properties [216–218]. Ferrofluids have been extensively studied and harnessed in a variety of applications [219–221]. For example,
the magnetorheological properties of magnetic colloids are exploited in highperformance bearings and seals [222–224]. There are several reviews on the properties and technological applications of ferrofluids [219–224].
A possibility of manipulating nanoparticles at a distance by applying an external
magnetic field opens up a broad field for biomedical applications of magnetic
nanoparticles. Specific surface groups of the stabilizers can be used for linking
different species to the nanoparticle surface. Drugs, antibodies, etc. can be attached
to the surface of colloidal magnetic nanoparticles and can be steered into regions
of the body where they are required by means of an applied magnetic field [203].
Functionalized magnetic nanoparticles can be specifically attached to a cancer
tumor and heated up by an a.c. magnetic field, resulting in thermoablation of the
tumor (magnetic fluid hyperthermia) [225]. Magnetic nanoparticles coupled with
oligonucleotides were used as a nanosensor for the detection and separation of
complementary specific oligonucleotides [226].
3.2 Synthesis of Metal Nanoparticles
The preparation of nanoparticles of desired sizes is the first and very important
step, being a prerequisite of their further investigation and use in practice. Narrow
size distribution or monodispersity is highly desirable because magnetic properties
become strongly size dependent in the nanometer size range. The magnetic behavior of nanoparticles depends not only on their chemical composition and size,
but also on their crystalline modification and the presence of structural defects like
stacking faults or twinned planes [210]. Murray et al. proposed the use of the term
‘‘nanocrystals’’ for crystalline particles with low concentrations of defects, while the
more general term ‘‘nanoparticle’’ also includes particles containing gross internal
grain boundaries, fractures, or internal disorder [227]. In the case of granular films
composed of magnetic nanoparticles, the interparticle spacing strongly affects the
magnetic and charge transport properties [228].
Since the different preparation methods lead to magnetic nanoparticles with
differences in crystalline structure, surface chemistry, shape, etc., the fabrication
technique has a great influence on the magnetic properties of the materials obtained [229], and, indeed, many examples are reported where differently prepared
nanocrystalline materials with similar grain sizes exhibit very different magnetic
properties [230]. Numerous physical and chemical methods have been employed to
produce magnetic nanostructures. Molecular beam epitaxy [231], chemical vapor deposition [232], normal incident pulsed laser deposition [233, 234], sputtering [235], and
electrodeposition [236, 237] are widely used for the preparation of thin magnetic
films. Separated nanometer-sized nanoparticles of Fe, Co and Ni can be prepared
by the inert gas condensation technique, where metal is evaporated at a very high
temperature (@1500 C) into a high purity gas (e.g., helium) [238, 239]. Upon collision with the inert gas, metal atoms lose their kinetic energy and condense on a
cold finger in the form of an ultrafine powder. A modification of this method is
based on the evaporation of metals by arc discharge into a circulating gas mixture
of H2 and Ar and is used for the large-scale preparation of 20–30 nm particles of
Fe, Ni and FeaNi alloys. Usually the nanoparticles synthesized by this approach are
amorphous, and some amount of oxide is present at their surface [238, 239]. An
arc discharge technique similar to that used in the fullerene synthesis was applied
for the preparation of carbon-encapsulated Ni, Co or Fe nanoparticles. Laser pyrolysis is a technique based on heating, with a laser beam, an organometallic reactant
vapor, e.g., a metal carbonyl mixed with an inert gas. Reactant vapors rapidly decompose to give the atoms, which form clusters upon collision with inert gas
molecules [229]. Nanoparticles of iron and iron carbides have been prepared in this
way. More detailed information about these methods can be found in [229].
Solution phase chemical syntheses are commonly used for preparation of colloids of magnetic metal nanoparticles. There are several general approaches used
for controlling the size and shape of nanoparticles in solution phase synthesis.
Thus, the water-in-oil microemulsion (reverse micelles) technique is based on using
droplets of water sustained in an organic phase by a surfactant. A metal salt dissolved in the nanodroplets can be reduced, inside these nanometer-sized reactors,
to metal nanoparticles. The amount of metal precursor available for the particle
growth is limited by the volume of the micelle, which can be varied by tailoring the
201
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3 Syntheses and Characterizations
water-to-surfactant ratio. This technique has been employed for the preparation of
nanoparticles of magnetic metals [240, 241], metal alloys [242], core-shell structures [243], and oxides [244, 245]. Thus, Pileni et al. synthesized nanoparticles of
Co, Fe or FeaCu alloy in water-in-isooctane microemulsion by the reduction of
metal salts using hydrazine or sodium borohydride in the presence of sodium
bis(2-ethylhexyl)sulfosuccinate, Na(AOT), as a surfactant [240]. The size of the
metal nanoparticles increased with increasing water-to-surfactant ratio ([H2 O]/
[AOT]) [246]. Synthesis in the reverse micelles is a fairly simple method and does
not require special costly equipment; however, several drawbacks of this approach
should be mentioned. First, the size distribution of nanoparticles is usually rather
broad (@21%) [247]. This means that monodisperse colloids can be obtained only
after costly post-preparative size-selective fractionation procedures [248]. Another
important problem is that nanoparticles of the magnetic metals do not crystallize
well at room temperature [249], and annealing at 200–300 C is required in order
to prepare nearly defect-free nanocrystals [250]. The reactions in the reverse micelles occur at relatively low temperatures (below 100 C), and the resulting material is usually poorly crystalline [251].
Sonochemical synthesis is based on the decomposition of metal precursor (usually
a metal carbonyl complexes) dissolved in a high-boiling solvent under irradiation
with high intensity ultrasound. Ultrasonic treatment initiates acoustic cavitation,
i.e. the formation, growth, and collapse of bubbles within a liquid. The collapse of
these bubbles locally heats up the media to thousands of kelvin. Volatile organometallic compounds rapidly decompose inside the collapsing bubbles, yielding the
individual metal atoms. In the presence of stabilizing agents these atoms agglomerate to nanosized particles [252]. This technique has been developed by Suslick et
al. for the preparation of Ni [253], nanostructured silica-supported Fe and FeaCo
alloys [254], and Fe [252] nanoparticles. Sonochemical synthesis usually yields
amorphous, rather polydisperse nanoparticles, and further high-temperature heating is required for their crystallization [252–254].
Among various solution phase approaches, the recently developed colloidal organometallic syntheses seem to be the most promising routes to the preparation of
high-quality magnetic nanocrystals. These methods are based on the reduction of
metal salts or the thermolysis of zero-valent organometallic precursors in the
medium of a high-boiling solvent in the presence of special stabilizing agents
which adsorb reversibly to the nanoparticle surface, mediating the growth rate. The
proper choice of surfactants, precursors, and temperature regime allows the successful preparation of monodisperse colloids of highly crystalline magnetic nanoparticles, as described in the next Section.
Organometallic Synthesis of Magnetic Nanoparticles
Most of the high-quality colloidal magnetic nanoparticles were synthesized via the
organometallic routes based on the thermolysis or reduction of organometallic
precursors in the presence of ligand molecules, so-called stabilizing agents which
bind to the nanoparticle surface, preventing its further growth and coagulation
with other particles. The proper choice of the stabilizing agents is of great impor3.2.2.2
3.2 Synthesis of Metal Nanoparticles
tance, as the attractive magnetic interactions between nanocrystals tend to decrease
the overall stability of the magnetic colloids. A high reaction temperature (150–
300 C) allows the annealing of crystalline defects and, consequently, is important
for the preparation of highly crystalline nanoparticles. In this Section we discuss
the main peculiarities of the organometallic synthetic approach using Co and
CoPt3 nanocrystals as the examples. Cobalt nanocrystals are probably the most
extensively studied nanocrystalline elemental magnetic materials, and CoPt3 can
be regarded as a model system for magnetic alloy nanocrystals.
General Remarks on the Synthesis of Co and CoPt3 Nanocrystals
The organometallic synthesis of magnetic metal nanocrystals usually requires a
rigorously air-free atmosphere and is carried out under a dry inert gas using
Schlenk line technique or a glovebox.
The size and size distribution of the nanoparticles are usually determined by the
kinetics of the particle nucleation and growth [255]. A narrow particle size distribution can be obtained if a temporally discrete nucleation event is followed by
controlled growth on the preformed nuclei [256]. Temporally discrete nucleation in
the organometallic synthesis can be attained by the hot injection technique schematically shown in Figure 3-92. A solution of organometallic precursors or a reducing agent (solution A) is rapidly injected into a hot solution containing stabilizing agents (solution B), resulting in explosive nucleation, and this is followed by
a temperature drop and slow growth of pre-formed nuclei. The optimized injection
temperature, duration of heating and the proper choice of the precursors and stabilizing agents are the prerequisites for a successful synthesis of magnetic nanoparticles with controllable size and shape and narrow size distribution.
If the synthesis does not yield nanoparticles of the desired monodispersity,
the size distribution of the as-prepared nanoparticles can be narrowed by a sizeselective precipitation procedure which exploits the difference in the solubility of
Ar
Ar
solution A
TC
20°C
TC
nucleation
growth
solution B
140-250°C
Fig. 3-92 Schematic representation of the hot
injection technique usually employed in the
organometallic synthesis of magnetic
nanocrystals. A thermocouple (TC) controls
the temperature in the reaction vessel. The
compositions of solutions A and B are detailed
in the text for each particular recipe.
203
204
3 Syntheses and Characterizations
smaller and larger particles [248, 257]. A typical example of a method of carrying
out the size-selective precipitation of a nanoparticle colloid is as follows. A sample
of as-prepared nanoparticles with a broad size distribution is dispersed in an appropriate solvent, and a non-solvent is added dropwise, gradually destabilizing the
colloidal dispersion until the solution becomes slightly turbid. The largest nanoparticles in the sample exhibit the greatest attractive van der Waals and magnetic
forces and tend to aggregate before the smaller particles. The aggregates consisting
of the largest nanoparticles can be isolated by centrifugation or filtration and redissolved in any appropriate solvent. A further portion of the non-solvent is added
to the supernatant to isolate the second size-selected fraction, and so on. For magnetic nanoparticles, the size-selective precipitation can also be performed by applying an external magnetic field, as the larger particles precipitate in the magnetic
field before the smaller ones [227]. These procedures repeated several times allow a
series of monodisperse fractions to be obtained with a near-Gaussian size distribution and a standard deviation below 5% (G1 atomic layer).
Synthesis of Cobalt Nanoparticles with Different Crystalline Modifications
Cobalt nanoparticles can be synthesized in three different crystalline modifications: hexagonal close-packed (hcp), face-centered cubic (fcc) and epsilon (e). The
e-Co structure, which is not observed in the bulk, consists of a twenty-atom unit
cell with cubic symmetry, similar to the structure of b-Mn [258]. The hcp structure
is the most stable phase for bulk cobalt at room temperature. Despite this fact,
experimental data repeatedly show that e-Co is the most often found crystalline
structure in nanoparticles prepared by wet chemistry [259].
The e-Co nanoparticles can be synthesized via reduction of a cobalt salt by superhydride (LiBEt3 H) in the presence of stabilizing agents [227]. Injection of the
solution of superhydride in dioctylether (solution A) into the hot (200 C) solution
of CoCl2 in dioctylether also containing oleic acid and trioctylphosphine (TOP) as
stabilizing agents (solution B) yields relatively monodisperse e-Co nanoparticles
(Figure 3-93a) [210]. HRTEM measurements (Figure 3-94b) and simulations of
X-ray powder data indicate that the e-Co nanoparticles are nearly perfect single
crystals [260]. The width of the XRD reflexes decreases coarsely with increase of
the nanoparticle size, as shown in Figure 3-93a for the (221), (310) and (311) reflexes. The average nanocrystal size can be tuned from @2 up to 6 nm by tailoring
the concentration of stabilizers. Increasing the concentration of oleic acid and TOP
yields a smaller mean particle size. In order to prepare e-Co nanocrystals larger
than 6 nm, one has to use less bulky stabilizers which can effectively cap particles
with lower surface curvature. Thus, substitution of tributylphosphine (TBP) for
TOP results in the formation of larger, 7–11 nm e-Co nanocrystals [227, 261]. Fine
tuning of the particle size can be performed by size fractionation of crude solutions, yielding a series of nearly monodisperse e-Co nanocrystals (Figure 3-94a).
Another method of preparation of cobalt nanoparticles is based on the thermal
decomposition of Co2 (CO)8 in the presence of surfactants [227]. A typical recipe
involves the injection of a solution of Co2 (CO)8 in diphenylether (solution A) into
the hot (200 C) mixture of oleic acid and TOP or TBP, also dissolved in diphenyl-
3.2 Synthesis of Metal Nanoparticles
Fig. 3-93 X-ray data (circles) for cobalt
nanoparticles and their corresponding
simulations (solid lines) using atomic-level
models of the particles. a e-Co particles,
b mt-fcc particles. The insets are X-ray data
and simulations for several particle sizes.
Reproduced from [260], Copyright 2001, with
permission from the American Chemical
Society.
ether (solution B). The injection is followed by heating at 200 C for 15–20
minutes. Decomposition of cobalt carbonyl under the conditions described above
results in the formation of Co nanoparticles with a so-called multi-twinned facecentered cubic (mt-fcc) lattice (Figures 6b and 7d). The multi-twinned particles are
composed of domains with a distorted fcc lattice. This structure is similar to that
observed in multiply-twinned, icosohedral gold particles [227]. Post-preparative
size selection procedures allow the isolation of fairly monodisperse fractions of
mt-fcc Co nanocrystals (Figure 3-94c). Depending on the reaction medium and
stabilizers, thermal decomposition of cobalt carbonyl can also yield nanoparticles
of other crystalline phases [259, 262–264].
205
206
3 Syntheses and Characterizations
Fig. 3-94 TEM images of cobalt nanoparticles.
a Low magnification image of e-Co. The even
contrast across the individual particles implies
a uniform crystalline structure for this phase.
b HRTEM image of e-Co showing perfect
crystallographic coherence and discrete
faceting. c Low-magnification image of the
mt-fcc structure. Contrast changes within
particles indicate they are composed of
crystal domains. d HRTEM image showing
the domains within a single mt-fcc particle.
These domains correspond to a multiply
twinned structure. Reproduced from [260],
Copyright 2001, with permission from the
American Chemical Society.
Annealing of e-Co nanocrystals at 300 C under vacuum converts them into the
hcp-Co phase (Figure 3-95), which has higher magnetocrystalline anisotropy than
that of e- and mt-fcc cobalt modifications [261]. Alternatively, the hcp-Co nanoparticles can be synthesized in solution by reduction of cobalt acetate by 1,2dodecanediol at 250 C in diphenylether medium in the presence of oleic acid and
TOP as stabilizing agents [227, 261].
As shown above, an optimized combination of precursors (CoCl2 , Co(CH3 COO)2
or Co2 (CO)8 ) with reducing agents (superhydride or polyalcohol) allows the selective preparation of monodisperse cobalt nanoparticles with a desired crystalline
phase. Such behavior shows that solution-phase chemical synthesis of magnetic
nanocrystals is not thermodynamically controlled and thus can enable the preparation of crystal phases which are metastable, like the e-Co structure [210]. The
control of size and crystalline phase of nanoparticles is important, because these
parameters strongly affect their magnetic properties, as discussed in Section
3.2.2.3 below.
3.2 Synthesis of Metal Nanoparticles
Fig. 3-95 XRD patterns of e-Co particles before and after
annealing: a 9 nm particles as synthesized and b after
annealing at 300 C under vacuum for 3 h. Reproduced from
[210], Copyright 1999, with permission from the American
Institute of Physics.
A serious drawback of cobalt and other elemental magnetic nanoparticles (Fe,
Ni, etc.) is their relatively low stability toward oxidation under ambient conditions.
Thus, the exposure of cobalt nanoparticles to air results in the formation of a shell
of @1–2 monolayers of cobalt oxide. This shell of anti-ferromagnetic cobalt oxide
considerably disturbs the magnetic properties of the cobalt core. Additionally, the
intrinsic coercivity of elemental cobalt is too low for Co nanoparticles to be considered
as candidates for advanced magnetic recording media. Nanoparticles of more stable materials with higher coercivity are highly desirable. Intermetallic materials such
as Cox Pt1x and Fex Pt1x seem to satisfy both the requirements mentioned above.
They are considerably more stable than elemental cobalt or iron and can have
crystalline modifications with very high magnetocrystalline anisotropy (e.g. L10 ).
The Synthesis of CoPt3 Magnetic Alloy Nanocrystals
The synthetic approach developed for the preparation of elemental nanoparticles
can be further extended to intermetallic compounds. Thus, high-quality CoPt3
nanocrystals can be synthesized via the simultaneous reduction of platinum acetylacetonate and thermal decomposition of cobalt carbonyl in the presence of 1-
207
208
3 Syntheses and Characterizations
adamantanecarboxylic acid (ACA) and hexadecylamine (HDA) as stabilizing agents
[265].
In a typical preparation, the mixture of platinum acetylacetonate, 1,2-hexadecanediol, 1-adamantanecarboxylic acid, and hexadecylamine (solution B in Figure 3-92) was heated to a certain temperature in the range 140–220 C, and solution A, prepared by dissolving cobalt carbonyl in 1,2-dichlorobenzene, was swiftly
injected into solution B with vigorous stirring. The temperature of injection of the
stock solution is henceforth referred to as the ‘‘reaction temperature’’. After injection, the color changes from pale yellowish to black, indicating the formation of
CoPt3 nanocrystals. Further heating was usually continued for 1 h at the injection
temperature and was followed by annealing at refluxing temperature (@275–
285 C) for 1 h. The annealing stage is necessary to improve the crystallinity of the
as-formed CoPt3 nanoparticles.
Information about nanoparticle crystallinity can be obtained by comparison of
the average particle sizes estimated by powder XRD and TEM methods. The width
of the XRD reflexes provides information about the X-ray coherence length, which
is close to the average size of the single crystalline domain inside the nanocrystal,
whereas TEM images show the total size of a nanoparticle. Nanocrystal sizes estimated by these methods are referred to as ‘‘XRD-size’’ and ‘‘TEM-size’’, respectively. For CoPt3 nanoparticles prepared at 140–200 C, the XRD-size is considerably smaller than the TEM-size (Figure 3-96a), because of the internal structural
defects (stacking faults, twinned planes, etc.) present within as-formed particles.
Annealing of CoPt3 nanoparticles at the boiling point of the crude solution (@275–
285 C) for @30 min results in an increase in the XRD-size, whereas no significant
change in the TEM-size is observed. The annealing process requires relatively high
temperatures (200–300 C) to trigger atom diffusion inside the nanocrystals. After
annealing, the XRD-size becomes almost equal to the TEM-size (Figure 3-96b), in-
Fig. 3-96 In situ annealing of CoPt3 nanoparticles. Comparison
of the XRD-size and TEM-size before and after annealing.
Reproduced from [255], copyright 2003, with permission from
the American Chemical Society.
3.2 Synthesis of Metal Nanoparticles
dicating that most of nanocrystals are single crystallites with a nearly perfect lattice. High-resolution TEM (HRTEM) investigations additionally confirm the excellent crystallinity of annealed CoPt3 nanocrystals (Figure 3-97). Almost all particles
possess lattice fringes without stacking faults or other defects. From a practical
point of view, perfect crystallinity of magnetic nanoparticles is important, because
the structural defects considerably reduce coercivity, the parameter determining
the applicability of a material for magnetic recording and storage [210].
Tuning of the size of CoPt3 nanoparticles can be achieved by control over nucleation and growth rates, as illustrated in Scheme 3-17. Fast nucleation provides a
Fig. 3-97 An HRTEM image of annealed CoPt3 nanocrystals.
Note that almost all nanoparticles possess the lattice fringes
without stacking faults or other defects.
nucleation
growth
fast
until all monomer is
consumed
precursors
slow
Schematic representation of the synthesis of
CoPt3 nanocrystals. Reproduced from [255], copyright 2003,
with permission from the American Chemical Society.
Scheme 3-17
nanocrystals
209
210
3 Syntheses and Characterizations
Fig. 3-98 a Dependence of CoPt3 nanocrystals size on the
reaction temperature. b Influence of initial amount of
stabilizing agent (1-adamantanecarboxylic acid) on the size
of the nanocrystals. Reproduced from [255], copyright 2003,
with permission from the American Chemical Society.
high concentration of nuclei, ultimately yielding smaller nanocrystals, whereas
slow nucleation provides a low concentration of seeds consuming the same
amount of precursors, thus resulting in larger particles [255].
The balance between nucleation and growth rates can be tuned by changing the
reaction temperature, because the activation energy for homogeneous nucleation is
usually much higher than that for particle growth [266]. This means that the
nucleation rate is much more sensitive to changes in temperature than the growth
rate. At higher reaction temperature, more nuclei are formed, and, according to the
scheme in Figure 3-96, the final particle size is smaller. Indeed, an increase in
temperature from 145 to 220 C allows the average size of the CoPt3 nanocrystal to
be reduced from @10 nm to 3 nm, i.e., it permits the preparation of nanocrystals
in a wide range of sizes (Figure 3-98a). This concept of size tuning implies that
Ostwald ripening, i.e. the growth of large particles with concomitant dissolution of
small ones, is very slow. As reported in [249, 255], this is indeed the case in the
most organometallic preparations of metal particles.
The nanoparticle size can also be tuned by tailoring the concentration of the
stabilizing agents. The effect of varying the concentration of stabilizers on the size
of the magnetic nanocrystals can be different for different systems. Thus, an increase in the concentration of the stabilizers resulted in smaller nanocrystals in the
case of Co [227, 267] and Ni [227], whereas an opposite behavior was observed for
iron [268] and CoPt3 [255] nanocrystals (Figure 3-98b).
For the kinetically driven synthesis of nanoparticles, a decrease in the mean
particle size with increasing stabilizer concentration can take place if strong passivation of the nanocrystal surface results in slowing the growth rate, whereas the
nucleation rate remains less influenced (Scheme 3-17). An opposite situation can
3.2 Synthesis of Metal Nanoparticles
Fig. 3-99 a Angular dependence of scattered
X-ray intensity (SAXS) from two samples of
as-prepared CoPt3 nanocrystals with mean size
(1) @4 nm and (2) @10 nm. b and d: Particle
volume fraction vs size curves calculated from
the SAXS data shown in part a. c and e: Particle
size distribution histograms estimated for the
same samples from TEM images. Reproduced
from [255], copyright 2003, with permission
from the American Chemical Society.
also be realized, as was recently observed for CoPt3 [255] and Fe-oxide [268] nanocrystals. Stabilizing agents like 1-adamantanecarboxylic acid (ACA) and primary
amines can also form stable complexes with individual metal atoms of a molecular
precursor. Thus, the platinum precursor (platinium acetylacetonate) decomposed
in the hexadecylamine (HDA)/diphenyl ether/1,2-hexadecandiol mixture at
@130 C, whereas addition of ACA increased the decomposition temperature to
@220 C. Similar behavior was observed for the cobalt precursor: addition of ACA
drastically enhanced the stability of the cobalt carbonyl solution toward thermal
decomposition. Formation of these complexes precedes the nucleation step [269]
and decreases the monomer reactivity. Therefore, an increase in the amount of
ligand would be expected to decreases nucleation rate. In accordance with Scheme
3-17, slowing the nucleation rate will result in a larger final size of the nanocrystals, and this was also observed experimentally (Figure 3-98).
Information about the in situ size distribution of CoPt3 nanoparticles can be
obtained by statistical evaluation of all nanocrystals in a representative TEM micrograph, or from small-angle X-ray scattering (SAXS) data by the indirect Fourier
transformation technique [270, 271]. The latter method provides the information
averaged over a huge number of nanoparticles and, moreover, allows the in situ
investigation of particle aggregation. Figure 3-99a shows the SAXS patterns for
two samples of as-prepared CoPt3 nanocrystals with mean sizes of @4 and 10 nm.
The nanocrystal size was tuned by varying the reaction temperature (215 C and
145 C). Figures 12b and c show the particle size distribution functions of asprepared @4 nm CoPt3 nanocrystals as obtained from the SAXS and TEM data,
211
212
3 Syntheses and Characterizations
respectively. The SAXS data provide evidence that small CoPt3 nanocrystals form
stable suspensions of isolated nanoparticles (single-particle population) [272, 273].
Both SAXS and TEM measurements confirm a relatively narrow and nearly symmetric particle size distribution for as-prepared @4 nm CoPt3 nanocrystals (Figures 12b and c), whereas a minor deviation between the mean sizes estimated by
these methods was observed. HRTEM always provided lower estimates of particle
size compared to SAXS, because the size distribution from SAXS is a volume (or
mass) distribution, whereas that from microscopy is a number distribution. The
difference between volume and number distribution is negligible when the distribution is narrow. However, in broad distributions or in the presence of a few aggregates, the larger particles count visibly more in volume distributions (SAXS)
than in number distributions (TEM).
In the case of @10 nm CoPt3 nanocrystals, the SAXS size distribution curve is
superimposed on a peak, corresponding to the single-particle population and a tail
in the large size region (Figure 3-99d). The TEM investigations of the same sample
revealed a symmetrical size distribution curve without any features around 12–13
nm (Figure 3-99e). This allows us to assume partial aggregation of @10 nm CoPt3
at room temperature [272]. The long-term stability of the colloidal solutions as well
as the absence of agglomerates in the TEM images provides evidence that the aggregation is reversible.
The absolute width of the particle size distribution is nearly constant for both the
@4 nm and the @10 nm samples (Figure 3-99). This means that the relative size
distribution of as-prepared @4 nm CoPt3 nanocrystals is much broader than that of
@10 nm nanocrystals (@14% and @7% of standard deviation, respectively). The
general tendency is a narrowing of the size distribution with increase in particle
size, as observed for all samples irrespective of the method of size control. Progressive narrowing of the particle size distribution with increasing in mean size is
due to the faster growth of the smaller particles compared with that of the larger
ones [266]. This growth regime was recently observed for ensembles of CdSe
nanocrystals and was called ‘‘focusing’’ of size distribution [274].
A combination of the methods of size control described above allows us to obtain
a series of nearly monodisperse CoPt3 nanocrystals ranging from @3 to 17 nm. It
was found that different experimental conditions (reaction temperature, ratio between the cobalt and platinum precursors, etc.) yielded nanocrystals of identical
chemical composition. Figures 3-100 and 3-101 show the systematic evolution of
XRD patterns and HRTEM images of CoPt3 nanocrystals with increase in their size.
CoPt3 nanocrystals with sizes smaller than @7 nm are usually spherical (Figures
3-97 and 3-101). However, further increase in the nanocrystal size results in a rather
abrupt transition from spherical to cubic, truncated cubic or, in some cases, platelike shapes, as shown in Figure 3-100. Annealing at @275 C results in smoothing
of the edges of cubic nanocrystals formed at 145 C (cf. Figures 3-101 and 3-102).
Each side of the cubic crystal lattices corresponds to the [100] direction of the fcc
nanocrystal lattice. Cubic particles can be formed if the growth rate in the [111]
direction is higher than that in the [100] direction [275].
Generally, shape anisotropy of magnetic nanocrystals can result in advanced
magnetic properties and is therefore of great practical interest [230]. The shape of a
3.2 Synthesis of Metal Nanoparticles
Fig. 3-100 Size-dependent XRD patterns of CoPt3 nanocrystals.
The average nanocrystal sizes were calculated using the DebyeSherrer equation. Reproduced from [255], copyright 2003, with
permission from the American Chemical Society.
colloidal nanoparticle can be controlled by selective adsorption of organic surfactants onto particular crystallographic facets in order to inhibit the particle growth
in a particular crystallographic direction. This approach to shape control is based
on the presence of chemically dissimilar lattice faces, e.g., the unique (001) face of
an hcp lattice. Shape-controlled colloidal syntheses were recently developed for
semiconductor (CdSe [276, 277], PbS [278], ZnO [279]) and noble metal (Pt [248],
Au [249]) nanoparticles.
A possibility for achieving shape control was also reported for Co nanocrystals
[259, 263, 264]. Thus, decomposition of cobalt carbonyl in the presence of primary
amines (e.g., octadecylamine) together with TOPO or oleic acid yields nanodisks of
hcp-cobalt (Figure 3-103) [264]. Alkylamines inhibit the growth of the unique (001)
face of hcp-Co upon increase of the particle size in a kinetic regime. Both length
and diameter of the nanodisks can be controlled by tailoring the reaction time
as well as by variation of the precursor-to-surfactant ratios. The long-term heating
of hcp-Co nanodisks (thermodynamic limit) results in the transformation of the
thermodynamically stable spherical nanoparticles of the e-Co phase. The nanodisks
are ferromagnetic and can spontaneously self-assemble into long ribbons formed
from stacked nanodisks lying perpendicular to the substrate. Note that a side view
of the nanodisks in transmission electron microscopy (TEM) images (Figure 3-103)
213
214
3 Syntheses and Characterizations
Fig. 3-101 TEM and HRTEM images showing
the effect of reaction temperature on the mean
size and size distribution of 3.7 nm, 4.9 nm,
6.3 nm, and 9.3 nm CoPt3 nanocrystals
prepared at 220, 200, 170, and 145 C. Molar
ratio of Pt:Co:ACA is 1:3:6 for all samples.
Reproduced from [255], copyright 2003, with
permission from the American Chemical
Society.
strongly resembles nanorods, and TEM tilting experiments are necessary to verify
the disk shape [264].
The synthesis of hcp-Co nanorods and nanowires was also reported [283]. Thermal decomposition of the organometallic cobalt precursor, [Co(h 3 -C8 H13 )(h 4 C8 H12 )], in the presence of oleic acid and oleylamine under a pressure of 3 bar H2
yielded @3 nm spherical hcp-Co nanoparticles, which were further transformed
into uniform nanorods (Figure 3-104) after heating for 48 h at 150 C. Both lengths
3.2 Synthesis of Metal Nanoparticles
Fig. 3-102 TEM overview image of faceted CoPt3 nanocrystals
prepared at 145 C. Shape-controlled synthesis of magnetic
nanoparticles.
Fig. 3-103 TEM image of partially self-assembled Co
nanodisks. Bar is 100 nm. Reproduced from [264], Copyright
2002, with permission from the American Chemical Society.
215
216
3 Syntheses and Characterizations
Fig. 3-104
TEM and HRTEM micrographs of cobalt nanorods. Courtesy of Dr. B. Chaudret.
and diameters of the hcp-Co nanorods can be varied by tailoring the nature and
concentration of the stabilizing agents. Thus, increasing the concentration of oleic
acid allows the preparation of very long (micron range) nanowires of diameter 4
nm. Conversely, replacing oleylamine by octylamine results in smaller and wider
Co nanoparticles (17 nm 10 nm). The mechanism of transformation of the initially formed @3 nm spherical nanoparticles into rods and wires is still unclear. In
principle, the nanorods could be obtained either through coalescence of spherical
particles followed by fusion [279–284] or by using the initial particles as nuclei for
an anisotropic growth process [285].
Colloidally synthesized CoPt3 nanocrystals possess chemically disordered facecentered cubic (fcc) structure which has no unique crystallographic direction.
Despite this, the formation of CoPt3 nanowires took place if the reaction was
performed at a relatively low temperature (@100–130 C) in the presence of 1adamantanecarboxylic acid and hexadecylamine as stabilizing agents [265]. The
CoPt3 nanowire-rich fraction (Figure 3-105a) can be isolated from the crude solution by centrifugation. The wires typically comprised two ‘‘heads’’, the ends of
which were approximately the same size as the spherical particles present in the
same sample, and a ‘‘body’’ with some bulges (Figure 3-105b). The concentration
of wires and their length decreased drastically upon prolonged heating of the
reaction mixture. After continuous heating at 280 C the quantity of nanowires
became negligible (Figure 3-105c).
Other Metal Magnetic Nanoparticles Synthesized by the Methods of
Colloid Chemistry
In the previous Section we described the colloidal synthesis of Co and CoPt3 nanocrystals and showed the great potential of the organometallic approach to the
preparation of high-quality magnetic nanoparticles. These synthetic schemes can
3.2 Synthesis of Metal Nanoparticles
Fig. 3-105 TEM images of CoPt3 nanocrystals
grown in a 1-adamantanecarboxylic acidhexadecylamine-diphenyl ether mixture: a
nanowire-rich CoPt3 fraction; b enlarged views
of nanowires; c spherical CoPt3 nanocrystals
formed from the nanowire-rich sample after
refluxing for 3 h. Reproduced from [265],
Copyright 2002, with permission from the
American Chemical Society.
also be adopted for the preparation of nanoparticles of different magnetic materials. In this Section, we briefly describe the preparation of some other important
magnetic nanoparticles.
Murray et al. reported the synthesis of monodisperse iron nanoparticles by the
thermal decomposition of Fe(CO)5 under an inert atmosphere in the presence of
oleic acid/tri-butyl phosphines [227]. These @6 nm Fe nanoparticles have a structure resembling the bulk bcc lattice with some disorder and are very sensitive to
oxidation. Even a short contact of a nanoparticle surface with air resulted in the
formation of oxide with a thickness of @2 nm [227]. Thermal decomposition of
217
218
3 Syntheses and Characterizations
iron carbonyl in tri-n-octylphosphine oxide at 320–340 C yielded @2 nm iron
nanoparticles [264]. Hyeon et al. used them as building blocks for the formation of
iron nanorods [286]. Thus, refluxing the 2 nm nanoparticles in pyridine in the
presence of didodecyldimethylammonium bromide (DDAB) resulted in the formation of uniform rod-shaped iron particles of diameter 2 nm and length varying
from 11 to 27 nm, depending on the concentration of the DDAB in solution in
pyridine. The electron diffraction pattern of the obtained nanorods exhibited a
body-centered cubic (bcc) structure of a-Fe. The transformation of nanospheres to
nanorods was explained by the oriented attachment of several 2 nm spherical particles followed by their fusion. The shape-dependent magnetic properties of iron
nanoparticles is discussed in the next Section. Practical applications of iron nanoparticles are impractical because of their chemical instability under ambient conditions. On the other hand, monodisperse iron particles can be gently oxidized by
trimethylamine N-oxide to form stable monodisperse magnetic g-Fe 2 O3 nanocrystals [268].
Spherical nickel nanoparticles have been synthesized by the reduction of nickel
salts by polyalcohols [227]. The mixture of oleic acid, tri-butylphosphine (TBP) and
tri-butylamine (TBA) as stabilizing agents favors the formation of @12–13 nm
nanoparticles. Increasing the length of alkyl chains in the stabilizing agents results in a decrease in the nanoparticle size. Thus, replacing TBP and TBA with
TOP and tri-octylamine (TOA) results in 8–10 nm Ni nanoparticles. Decomposition of Ni(cycloocta-1,5-diene)2 in the presence of HDA as the stabilizing and
shape-controlling agent yielded Ni nanorods [287]. Simultaneous reduction of
nickel and cobalt salt was used to prepare cobalt-nickel alloy nanoparticles with the
fcc structure [227]. Nickel is more readily incorporated into the growing nanoparticles, and so employing equal amounts of nickel and cobalt precursors leads to
the formation of an Ni-rich alloy (@Ni60 Co40 ).
Some of the most promising magnetic nanoparticle materials, which may give a
boost to data storage technology, are the iron-platinum alloys [288]. The classical
synthesis by Sun et al. is based on the simultaneous high-temperature reduction of
platinum acetylacetonate and decomposition of iron pentacarbonyl in the presence
of oleic acid and oleyl amine as stabilizing agents [289]. The size of the PtFe
nanocrystals can be controlled by first growing @2.5–3 nm monodisperse seed
particles and then adding more reagents to enlarge the existing seeds until the desired size, @10 nm, is reached [289]. The chemical composition of Fex Pt1x binary
alloy nanocrystals can be tuned between x ¼ 0:48 and x ¼ 0:7 by controlling the
molar ratio of iron carbonyl to platinum salt.
As-prepared FePt nanoparticles exhibit the chemically disordered fcc structure.
fcc FePt nanoparticles are superparamagnetic at room temperature and can be
handled in the form of a colloidal solution for casting films, making self-assembled
arrays, etc. (Figure 3-106). Further, the FePt nanoparticles can be annealed at 550–
600 C under an inert atmosphere. This annealing induces the iron and platinum
atoms to rearrange, converting the particles to the chemically ordered face-centred
tetragonal (fct) phase (Figure 3-107). Room temperature coercivity of the arrays of
fct FePt nanocrystals is high enough to use them in a data storage device [289].
3.2 Synthesis of Metal Nanoparticles
Fig. 3-106
Hexagonal close-packed 2D array of FePt nanocrystals on a TEM grid.
Fig. 3-107 XRD patterns (A) of as-synthesized
4-nm Fe52 Pt48 particle assemblies and a series
of similar assemblies annealed under atmospheric N2 gas for 30 min at temperatures
of (B) 450 C, (C) 500 C, (D) 550 C, and
(E) 600 C. The indexing is based on tabulated
fct FePt reflections (25). Reproduced from
[289], Copyright 2000, with permission.
219
220
3 Syntheses and Characterizations
Depending on the Fex Pt1x stoichometry and nanocrystal size, the transition
temperature was in the range 500–700 C [290], while for the bulk material the
reported value was 1300 C [291]. In order to further reduce the phase transition
temperature, it was proposed to introduce silver atoms into the lattice of FePt
nanocrystals during their synthesis [292]. Incorporation of Ag promotes the transition from fcc to the tetragonal phase, reducing the transition temperature by
@100–150 C compared to the pure FePt nanocrystals. It was suggested that Ag
atoms leave the FePt lattice at temperature of less than 400 C, and that the vacancies formed increase the mobility of the Fe and Pt atoms, accelerating the phase
transformation.
The synthetic approach developed for the synthesis of platinum-iron binary alloy
nanoparticles was adopted for the preparation of several other binary alloy nanoparticles, such as FePd [293] or MnPt [294]. However, these alloys require further
structural characterization as well as further development of the methods of control of particle size, shape and composition.
The possibility of preparing ‘‘core-shell’’-type magnetic nanoparticles was reported [295, 296]. These core-shell nanoparticles can be prepared by a two-step
synthesis in which nanoparticles of one metal serve as seeds for the growth of a
shell from another metal. Thus, reduction of platinum salts in the presence of the
Co nanoparticles enabled the preparation of air-stable Co-core/Pt-shell nanoparticles [295]. Thermal decomposition of cobalt carbonyl in the presence of silver salt
led to the formation of Agcore Coshell nanoparticles [296].
Summarizing, a combination of organometallic and colloid chemistry gives a
powerful method of preparation of monodisperse magnetic nanoparticles with
controllable size, shape, composition, and crystalline modification. The present
state of the art in the synthesis of magnetic nanoparticles allows us to provide a
palette of high-quality materials for investigations of nanomagnetism as well as for
various applications. In spite of the progress in the synthesis of magnetic nanocrystals, there is still a lack of theoretical understanding of the processes occurring
during the growth of nanoparticles in colloidal solutions and of the problem of
keeping the particle size distribution narrow. Further synthetic progress will definitely depend on our ability to understand and control the parameters governing
the properties of colloidally grown nanocrystals.
Size- and Shape-Dependent Magnetic Properties of Ferromagnetic Metal
Nanoparticles
Magnetism is a result of moving charges. Additionally, elementary particles, e.g.,
electrons, have an intrinsic magnetic moment (spin), which determine their
quantum state. The magnetic properties of materials arise mainly from the motion
of electrons: the orbital motion and the spin motion. Other contributions such as
that from nuclear magnetic effects are usually much smaller than that from the
electron. The electronic structure of a nanometer-sized metal particle is strongly
size dependent (see, e.g., Chapter 5.2), resulting in size-dependent magnetic behavior. Many physical phenomena (magnetic domain size, exchange coupling effects, etc.), which determine experimentally observable magnetic properties of ma3.2.2.3
3.2 Synthesis of Metal Nanoparticles
M
Ms
Mr
Hc
H
Fig. 3-108 A typical hysteresis loop for a ferromagnetic
material and the important magnetic parameters which can
be obtained from the hysteresis loop. Dashed line shows the
first scan.
terials, have natural length scales lying in the nano- and micrometer size range.
What happens if the size of a ferromagnetic material is shrunk below 10 nm? In
this Section we intend to show the effect of particle size, shape, and structure on
magnetic properties of nanomaterials. The reader can find a more detailed description of the physical aspects of nanoscale magnetism in several excellent books
[203, 229, 297] and journal reviews [227, 230].
First, we would like to very briefly introduce the reader to the commonly measured magnetic parameters. All materials interact with a magnetic field, and this
interaction can be either attractive toward a magnetic pole (ferro- and paramagnetism) or repulsive (diamagnetism). The application of a magnetic field (H) results in magnetization (M) of a sample, which can be measured by, e.g., a superconductive quantum interference device (SQUID), one of the most popular and
sensitive methods of investigating magnetic properties. When a ferromagnetic
material is magnetized by an increasing applied field and then the field is decreased, the magnetization does not follow the initial magnetization curve obtained
during the increase. This irreversibility is called hysteresis [203]. A typical hysteresis loop arises from measuring the magnetization of the material as a function of
magnetic field applied in positive and negative directions, i.e., the response of the
material follows two distinct paths on magnetization and demagnetization (Figure
3-108). At large fields, the magnetization approaches the maximum value, called
the saturation magnetization (Ms ). Magnetic materials in a ferromagnetic state
have a residual magnetization at zero external field, called the remanent magnetization (Mr ). Coercivity (Hc ) characterizes the reverse-field strength needed to reduce the magnetization to zero. Thus, hysteresis measurements allow us to obtain
information about coercivity, remanent magnetization, and saturation magnetization of a given material (Figure 3-108).
221
222
3 Syntheses and Characterizations
Tab. 3-2
Magnetic units and unit conversions.
Magnetic units
Magnetic field strength
Magnetization
Magnetic flux density
(induction)
Induction constant
H
M
B
m0
cgs
SI
Conversion
Oe (oersted)
emu/cm 3
G (gauss)
B ¼ H þ 4pM
G/Oe
A/m
A/m
T (tesla)
B ¼ m 0 ðH þ MÞ
T/(A/m)
1 Oe ¼ 10 3 =4p A/m
1 emu/cm 3 ¼ 10 3 A/m
1 T ¼ 10 4 G
m0 ¼ 1 G/Oe
4p 107 T/(A/m)
It is still common to use the old cgs (centimetre, gram, second) system for
magnetic units in scientific and engineering literature. As a consequence, there are
two widely used systems of units, the cgs and the SI system. Thus, the magnetic
field strength H can be measured in Oe (cgs) or in A/m (SI). Moreover, the applied
field can be multiplied by the induction constant m 0 and given in units of Gauss
(cgs) or Tesla (SI). The magnetization M is measured in emu/cm 3 (cgs) or A/m
(SI) or, being multiplied by m 0, can also be given in Gauss or Tesla. Also note that
in the cgs system the magnetization can also be written per gram of substance.
Table 3-2 presents the magnetic units used in cgs and SI systems and some unit
conversions.
To minimize their energy, macroscopic ferromagnetic materials are divided up
into domains of parallel magnetic moments. Within a domain, the magnetic moments orient in one direction, while the alignment of spins in neighboring
domains is usually antiparallel. The oppositely aligned magnetic domains are
separated from each other by a domain wall (Bloch wall). As the particle size decreases below some critical value, the formation of domain walls become energetically unfavorable, and the ferromagnetic particle can support only a single domain. This critical size depends on the material and is usually on the order of
tens of nanometers, varying from @14 nm for Fe up to @170 nm for g-Fe 2 O3 . One
can see that magnetic particles of nanometer size are usually a single domain
[227].
The dimensionless parameter – magnetic susceptibility (w) defined as w ¼
qM=qH is used to estimate the efficiency of the applied magnetic field for magnetizing a material. In many cases, the susceptibility of a material depends on the
direction in which it is measured. Such a situation is called magnetic anisotropy.
The physical origin of magnetic anisotropy can be, e.g., the symmetry of the crystalline lattice or the shape of a particular piece of magnetic material, as is considered below in more detail. Now we would only like to emphasize that the
anisotropy is an extremely important parameter determining the behavior of a
ferromagnetic material. When magnetic anisotropy exists, the total magnetization
of a system will prefer to lie along a special direction, called the easy axis of magnetization. For a small single-domain particle, the energy associated with this
alignment is called the anisotropy energy and can be written in the simplest
uniaxial approximation as Ea ¼ KV sin 2 y, where K is the anisotropy constant, V
Energy
3.2 Synthesis of Metal Nanoparticles
DE
0
p
p
q
2
Orientation
Fig. 3-109 The energy associated with rotation of a
magnetization direction for a uniaxial magnetic particle. The
minima correspond to alignment of the magnetization along
the easy axis (schematically shown by dotted lines). DE
represents the energy barrier to the rotation of magnetization.
the particle volume, and y the angle between the moment and the easy axis (Figure
3-109) [227].
When an external magnetic field interacts with the single-domain particle, an
~s H
~ is supplied. Depending on the direcadditional potential energy of Ef ¼ M
tion of the magnetic field with respect to the particle easy axis, the response of the
particle moment to an applied field is different. Thus, if the field is applied perpendicular to the easy axis, the dependence of the particle potential energy on the
direction of magnetization varies with the strength of applied field as shown in
Figure 3-110a. The equilibrium direction of the magnetic moment, corresponding
to the minimum of the total alignment energy E tot ¼ Ea þ Ef , turns toward the
field to the angle y ¼ arcsinðMs H=ð2KVÞÞ. At H ¼ 2KV=Ms the moment aligns
perpendicular to the easy axis (i.e. along the applied field) and magnetization approaches saturation. In the M vs. H scans, the magnetization component parallel
to the applied field changes linearly from Ms to þMs as shown in Figure 3-110c.
Note that no magnetization hysteresis is observed if the magnetic field is applied
perpendicular to the easy axis of a uniaxial single-domain particle.
If a magnetic field is applied parallel to the easy axis, the particle moment can
align either parallel or antiparallel to the field direction. The latter case is considered in Figure 3-110b. The alignment of the particle magnetization along the field
is energetically favorable; however, in a relatively weak field the energy barrier
blocks the magnetization reversal. The height of the barrier decreases with increasing field strength (Figure 3-110b). At H ¼ 2KV=Ms the barrier vanishes and
the particle magnetic moment jumps from antiparallel to parallel alignment. This
behavior leads to a square hysteresis loop with coercivity Hc ¼ 2KV=Ms as shown
in Figure 3-110d. The effect of magnetic field applied at any other angle to the
particle easy axis can be evaluated in a similar manner.
223
224
3 Syntheses and Characterizations
a
magnetic field
magnetic field
H
b
2KV
Ms
no field
c
M
M
Ms
d
Ms
2KV
Ms
2KV
Ms
H
a,b Dependence of potential energy
on the direction of a particle magnetic moment
in the presence of external magnetic fields
applied (a) perpendicular and (b) parallel to
the easy magnetization axis. The lower curves
correspond to zero magnetic field and the
upper ones correspond to the saturation
Fig. 3-110
H
magnetization when the magnetic moments
align along the field. c and d: Hysteresis loops
for applied fields perpendicular (c) and parallel
(d) to the easy axis of a uniaxial single-domain
particle. The direction of the axis of easy
magnetization is shown by dotted lines.
As one can see, the response of an individual single-domain particle crucially
depends on the direction of the applied magnetic field. The two situations considered above represent the extreme cases of the possible hysteresis curves, totally
closed (no hysteresis) and totally open (square). Other orientations yield hysteresis
curves between these limits. In reality, an ensemble of many particles has ran-
3.2 Synthesis of Metal Nanoparticles
strong field
no field
a
fixed particles
Fig. 3-111 a An ensemble of magnetic
particles with randomly distributed axes of
easy magnetization. b Magnetization of a
particle ensemble in a strong field via rotation
of the particle magnetic moments. c Magnetic
b
c
free particles
particles dispersed in a liquid can align their
easy axes along the applied field. Directions of
particle easy magnetization axes are shown by
dotted lines.
domly oriented easy axes (Figure 3-111a), and the measured hysteresis loop is the
result of averaging over all possible orientations. In an extremely strong magnetic
field all particles align their moments along the field, approaching the largest (saturation) magnetization. If the particles are fixed, e.g., dispersed in a solid nonmagnetic matrix, the easy axes are randomly aligned and the magnetization occurs
through rotation of the particle moments (Figure 3-111b). The time scale of this
process is approx. 109 s. However, the particles dispersed in a liquid can considerably decrease the potential energy aligning their easy axes along the external field
(Figure 3-111c). This alignment occurs via the Brownian motion process in which
the particles’ collision rate with solvent molecules is much slower (@107 s for
solvents like hexane) than the rate of rotation of magnetic moment inside the
particle.
Let us now consider the role of thermal fluctuations in the behavior of very small
magnetic particles. In zero magnetic field, the energy barrier DE has to be overcome in order to rotate the magnetization of a single-domain particle (Figure 3109). The height of this barrier DE ¼ KV is proportional to the particle volume,
and, as the particle size decreases, DE may become comparable to the thermal energy (kB T, where kB is the Boltzmann’s constant). In this case, the energy barrier
can no longer pin the direction of magnetization to the time scale of observation,
and rotation of the direction of magnetization occurs due to thermal fluctuations.
Such a particle is said to be superparamagnetic. The coercivity of a superparamagnetic particle is zero (Hc ¼ 0), because thermal fluctuations prevent the existence
of a stable magnetization. Cooling of a superparamagnetic particle reduces the
energy of thermal fluctuations, and, at a certain temperature, the free movement of
magnetization becomes blocked by anisotropy. The temperature of the transition
from the superparamagnetic to the ferromagnetic state is called the blocking temperature (TB ). The blocking temperature is related to the particle volume and the
anisotropy constant. The latter can be calculated as K ¼ 25kB TB =V [203].
The coercivity of a magnetic particle strongly depends on its size. Thus, if the
particle is large enough to support a multidomain structure, magnetization rever-
225
226
3 Syntheses and Characterizations
Fig. 3-112 Particle coercivity vs size. The largest coercivity is
observed at the particle size dc corresponding to the transition
from multidomain to single-domain structure.
sal occurs through a domain walls motion [230]. This is relatively easy, and hence
the coercivity is low. On the other hand, in a single-domain particle the change of
direction of magnetization can occur only by coherent rotation of spins, resulting
in considerably higher coercivity of single-domain particles than that of multidomain particles (Figure 3-112) [203]. Upon further decrease of the particle size,
the coercivity falls off because of the progressively increasing role of thermal fluctuations, leading to superparamagnetism, with Hc ¼ 0 (Figure 3-112).
In general, a magnetic nanoparticle is a single-domain magnet. To investigate
the magnetic properties inherent in independent nanoparticles, the measurements
have to be performed on an ensemble of uniform non-interacting nanoparticles.
All sources of sample inhomogeneity (particle size and shape distribution, presence of particles with different crystalline modifications, etc.) must be minimized.
Interparticle interactions can be reduced by dissolving the particles in an appropriate solvent at a low concentration [227]. If a distribution of particle sizes is
present, the initial susceptibility is more sensitive to the larger particles present,
whereas the approach to saturation is sensitive to the smallest particles of the
sample [230]. In addition to hysteresis measurements, which allow the coercivity
and specific magnetization of a sample to be estimated, zero-field-cooled (ZFC)
and field-cooled (FC) measurements provide information about the blocking temperature (TB ). In ZFC scans, a sample is cooled under zero applied magnetic field
to a temperature well below the suspected TB . The system is then warmed up, and
the magnetization is measured as a function of temperature by applying a relatively low external magnetic field. As the thermal energy increases, the nanoparticles align with the applied field and the magnetization increases. At the
blocking temperature the magnetization is maximal. Further increase in temperature above TB results in decrease in the magnetization because the thermal energy
causes fluctuation of magnetic moments of nanoparticles. In the field-cooled (FC)
scans, a sample is cooled in a small magnetic field, freezing-in a net alignment of
the nanoparticle moments. The field is then removed, and the magnetization is
measured as the sample is slowly warmed up. Thermal energy unpins and randomizes the nanoparticle moments, lowering the sample’s net magnetization. Be-
3.2 Synthesis of Metal Nanoparticles
low TB the free movement of magnetic moments is ‘‘blocked’’ by the anisotropy, so
that the particles are in a ferromagnetic state, whereas above TB the particles are
characterized by superparamagnetic behavior [247]. At the blocking temperature
the ZFC and FC curves converge [227]. The blocking temperature in a superparamagnetic system decreases with increasing applied measuring fields, and is
proportional to H 2=3 at large magnetic fields and to H 2 at lower fields [230].
Figure 3-112 shows a set of measurements performed for the characterization of
magnetic properties of cobalt nanoparticles [227]. As was discussed in Section
3.2.2.2, size series of monodisperse cobalt nanoparticles can be synthesized with
three different intrinsic crystal structures: hexagonal (hcp), e, and multi-twinned
face-centred cubic (mt-fcc). All investigated sizes are far below the critical single
domain size which is, e.g., 70 nm for the hcp cobalt phase, and the particles are
small enough to be superparamagnetic at room temperature. The hysteresis
measurements performed at 5 K on hcp-Co nanoparticles of different sizes revealed a strong decrease in coercivity with decreasing particle size (Figures 26a and
26d). As already discussed, the number of spins coupled by exchange interaction
decreases with the size of the single-domain particle, leading to an easier spontaneous reorientation of magnetization at a given temperature and thus to a decrease
in coercivity [227].
Reducing the size of a single-domain particle decreases the number of spins
exchange-coupled to resist spontaneous reorientation of its magnetization at a
given temperature [227]. The drop in saturation magnetization with decreasing
nanoparticle size (Figure 3-113a) was attributed to increase in nanoparticle surfaceto-volume ratio [227]. The measurements of the hysteresis loop at different temperatures show the transition from ferromagnetic to superparamagnetic behavior
in 9.5 nm e-Co nanocrystals (Figure 3-113b). Increasing thermal energy results in
fluctuations of direction of magnetization of the nanoparticles, reducing coercivity
and leading to superparamagnetic behavior at temperatures above 125 K. More
precise information about the blocking temperature can be obtained from ZCF and
FC scans, shown in Figure 3-110c for a size series of hcp-Co nanocrystals. The
blocking temperature estimated from the peak magnetization on the ZCF scans
increases with the particle size from 20 K for 3 nm up to 240 K for 11 nm hcp-Co
nanoparticles. The increase in blocking temperature, magnetization, and coercivity
with increase in nanoparticle size was observed for all crystalline modifications.
Similar behavior was also observed for CoPt3 nanocrystals [255].
The magnetic anisotropy of nanocrystalline materials originates from the influence of crystal symmetry, shape, and surface (or grain boundaries) effects. The
most common anisotropies in the case of magnetic nanoparticles are magnetocrystalline and shape anisotropies.
Magnetocrystalline anisotropy arises from the coupling of the electron spins to the
electronic orbit, which, in turn, is coupled to the crystal lattice. The magnetization
is energetically favorable to alignment along a specific crystallographic direction
called the easy axis of the material. In the case of hexagonal cobalt, the easy axis of
magnetization is the c axis. Magnetocrystalline anisotropy is intrinsic to a given
material and independent of particle shape [230] but dependent on temperature.
227
228
3 Syntheses and Characterizations
Fig. 3-113 a Magnetization versus appliedfield hysteresis loops at 5 K for 3, 6, 8, and
11 nm hcp cobalt nanocrystals. b Comparison
of the hysteresis loops for 9.5 nm e-Co nanocrystals measured at 5, 40, 100, and 300 K.
c A series of zero-field-cooled (ZFC) and field-
cooled (FC) magnetization scans for hcp-Co
nanocrystals of different sizes (3, 6, 8, and
11 nm). d Dependence of coercivity on
nanoparticle size for a collection of hcp-Co,
e-Co, and mt-fcc Co samples. Reproduced from
[227], Copyright 2001, with permission.
In a cubic lattice, symmetry creates multiple easy axes. At room temperature, the
easy magnetization axes of cubic Ni and Fe are h111i and h100i axes, respectively
[298].
The highest magnetization and coercivity were achieved in the case of hcp cobalt
nanoparticles (Figure 3-113d), because this structure has greater magnetocrystalline anisotropy than that of the mt-fcc and e-Co phases [227]. For each system, coercivity is lower than that predicted for idealized nanocrystals, as internal structural
defects (stacking faults, twinned planes) reduce Hc [210]. Future ultrahigh-density
3.2 Synthesis of Metal Nanoparticles
recording media will require uniform particles (size distribution less than 10%)
with an average diameter of <6 nm and a room temperature Hc of 3000–5000 Oe
[227]. The coercivity is proportional to the anisotropy constant, so monodisperse
nanoparticles of materials with high intrinsic magnetocrystalline anisotropy (e.g.,
platinum-iron and cobalt-platinum binary alloys) are essential [230].
Cobalt nanoparticles have a strong tendency to become oxidized in air, and their
surface is usually covered with a layer of anti-ferromagnetic cobalt oxide, which
influences the magnetic properties [230]. The influence of this oxide layer is especially noticeable in the case of the small nanoparticles with an extremely high
surface-to-volume ratio. Thus, the coercivity of the smallest nanoparticles in the
size series shown in Figure 3-110d is dominated by coupling to the surface CoO
layer [227, 230]. The phenomenon is called exchange anisotropy [299] and occurs
when a ferromagnet is in close proximity to an antiferromagnet or ferrimagnet
[230]. Exchange anisotropy arises as a consequence of the interfacial, exchange
coupling. This effect was discovered in oxide-coated cobalt particles (with sizes in
the range 10–100 nm), and since then it has been observed in a variety of other
systems.
Some molecular species adsorbed on the surface of nanoparticles can also significantly influence the magnetic behavior through the quenching of the surface
atom contribution. Complete quenching of the magnetic moment of the surface
nickel atoms was reported for small nickel particles and Ni38 Pt6 clusters coated
with carbonyl ligands [217].
Non-spherical nanoparticles can also possess shape anisotropy. Shape anisotropy
accounts for preferential orientation of the magnetization along the long axis of an
elongated particle. The demagnetization field is less in the long direction, because
the induced poles at the surface are further apart [203]. Shape anisotropy is comparable to magnetocrystalline anisotropy and can produce large coercive forces.
Thus, the blocking temperature of Fe nanorods (2 nm 11 nm) is nearly one
order of magnitude higher than that of 2 nm spherical nanoparticles (Figure 3-114)
[286]. The advanced magnetic properties of non-spherical nanoparticles has led to
great attention being paid to shape-controlled colloidal synthesis. The syntheses of
Co nanodisks [264], Co nanorods [283], Fe nanorods, CoPt3 nanocubes [255], and
CoPt3 nanowires [265] have been recently reported. Progress in the synthesis of
monodisperse non-spherical magnetic particles is the most important prerequisite
for the investigation of shape-dependent magnetic properties.
If the concentration of nanoparticles in a sample is high, magnetic moments of
individual nanoparticles can interact, and these dipolar interactions strongly affect
the magnetic behavior of the sample [300]. Interparticle interactions in 2D and 3D
assemblies of magnetic nanoparticles result in the shift of blocking temperature to
higher values. Thus, a 5 K increase in the Tb was observed for 5.8 nm cobalt particles when they were arranged in a 2D network instead of being isolated noninteracting particles [247]. Also, an increase in the blocking temperature and
broadening of the ZFC/FC scans was found in close-packed assemblies of monodisperse nickel [227] and CoPt3 nanoparticles. Dilution of a close-packed film of 9
nm cobalt particles (Tb ¼ 165 K) with 1-octadecylamine reduced the magnetostatic
229
230
3 Syntheses and Characterizations
3
Fig. 3-114 Zero-field-cooled (ZFC) and field-cooled (FC)
magnetization scans for 2 nm spherical iron nanoparticles and
2 nm 11 nm iron nanorods at an applied magnetic field of
100 Oe. Reproduced from [286], Copyright 2000, with
permission from the American Chemical Society.
coupling of the particles and resulted in a sharper transition from the superparamagnetic to the ferromagnetic state at 105 K [210].
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Z. L. Wang (ed.), Wiley-VCH,
Weinheim, 2000, p. 289.
298 C. Kittel, Introduction to solid state
physics, John Wiley & Sons, New York,
Singapore, 1996.
299 W. H. Meiklejohn, C. P. Bean, Phys.
Rev. 1957, 105, 904.
300 G. A. Held, G. Grinstein, H. Doyle,
S. Sun, C. B. Murray, Phys. Rev. B
2001, 64, 012408.
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4
Organization of Nanoparticles
4.1
Semiconductor Nanoparticles
Alexander Eychmüller
Once formed in solution or as powders, semiconductor nanocrystals may become
building blocks for larger supramolecular units. These units may form ‘‘automatically’’, i.e. by crystallization or self-assembly, or they may be formed ‘‘intentionally’’ using one of several techniques, and these methods of formation are discussed here. Structures of semiconductor nanoparticles as large as several
hundreds of microns can be formed as crystals (Section 4.1.1) or superlattices
(Section 4.1.2), depending on fine adjustment of the experimental conditions. Section 4.1.3 deals with a number of experimental techniques, e.g., the LangmuirBlodgett and layer-by-layer techniques, which have in common the build-up of
layered structures consisting of disordered particles, while Section 4.1.4 is devoted
to the coupling (either ionic or covalent) of semiconductor particles.
4.1.1
Molecular Crystals
One of the very interesting developments of the last decade is the merging of the
colloidal preparation of nanostructures with synthetic approaches that stem from
advanced inorganic chemical routes (see, e.g., Section 3.1.3), both leading to very
similar nanoparticles. Both routes yield species that crystallize into superstructures
with sizes reaching some hundreds of microns. Single-crystal X-ray diffraction
is performed on these macroscopic crystals, yielding the exact positions of every
single atom in the superstructures.
Based on the works of Strickler [1] and Dance and coworkers [2–5] in 1993,
Herron et al. reported on the crystallization of [Cd32 S14 (SPh)36 ]4DMF (Ph ¼
C6 H5 , DMF ¼ dimethyl formamide) as pale yellow cubes [6]. The structure unraveled by single-crystal X-ray diffraction consists of an 82-atom CdS core that is a
piece of the bulk cubic lattice appearing as a tetrahedron whose points are capped
by DMF solvent molecules. The cluster remains soluble, and as such retains properties of a discrete molecule. From the point of view of a nanocrystal chemist inNanoparticles: From Theory to Application. Edited by Günter Schmid
Copyright 8 2004 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
ISBN: 3-527-30507-6
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4 Organization of Nanoparticles
Fig. 4-1
Structures of two covalently linked [Cd17 S4 (SCH2 CH2 OH)26 ] clusters [7].
terested in studying the properties of particles as a function of size, the ‘‘Cd32 cluster’’ produced is exceptional because, unlike all other colloidally prepared
samples, it has no size distribution. Its structure is a larger homolog of the
[Cd17 S4 (SPh)28 ] 2 cluster prepared by Lee et al. [5]. This compound is obtained by
recrystallization from hot CH3 CN of the mixed products of the reaction of PhSH,
Et3 N, Cd 2þ , Na2 S, and Me4 NCl in methanol/acetonitrile. As in the case of the
‘‘Cd32 -cluster’’, this ‘‘Cd17 -cluster’’ appears as a noncovalently bound cluster in the
crystal. An identical cluster core was prepared by Vossmeyer et al. but via a distinctly different route and with another superstructure in the evolving crystals [7].
The preparative route was essentially described earlier as a colloid-chemical one in
which Cd(ClO4 )2 and 2-mercaptoethanol were dissolved in water, followed by H2 S
injection in alkaline solution [8]. Exhaustive dialysis against water yields crystalline
material which first appears as needles and evolves as blocky crystals after standing in mother liquor for some weeks. Figure 4-1 displays the structures of two
[Cd17 S4 (SCH2 CH2 OH)26 ] clusters that form the basic units of the superlattice. In
contrast to the above-mentioned clusters, here intermolecular bridges are formed
that build up a superlattice framework.
Smaller thiophenolate-capped clusters containing zinc, e.g. [Zn 4 (m2 -SPh)6 (SPh)4 ] 2 , [Zn10 (m3 -S)4 (m2 -SPh)12 ], and [Zn10 (m3 -S)4 (m2 -SPh)12 (SPh)4 ] 4 , are prepared as supertetrahedral fragments and possible molecular models for cubic ZnS,
and their electronic structure is evaluated theoretically [9].
As mentioned above, very similar structures are obtained via organometallic
synthetic routes. Clusters derived in this manner include [Cd32 Se14 (SePh)36 (PPh3 )4 ] [10] and its mercury analog, which display exactly the same internal
structure as the above-mentioned sulfur-containing clusters. Other crystalline material consists of [Cd17 Se4 (SePh)24 (PPh2 Pr)4 ] 2þ [Cd8 Se(SePh)12 Cl 4 ] 2 [11] (Pr ¼ npropyl) and the smaller cluster [Cd10 Se4 (SPh)12 (PPr3 )4 ] [12]. The latter work also
contains results concerning the cluster compound [Cd16 (SePh)32 (PPh3 )2 ], which
lacks a central atom, in contrast to the ‘‘Cd17’’ cluster, thus forming a structure
similar to that of Koch pyramids [13].
4.1 Semiconductor Nanoparticles
Together with the Banin group, the authors of [10–12] performed optical spectroscopy on some of the cluster molecules gained [14–16]. They were treated as the
molecular limit of the bulk semiconductor CdSe, and issues like the oscillator
strength, the steady-state and time-resolved photoluminescence and the photoluminescence excitation were addressed. Additionally, emission-mediating vibrational modes were detected, and photobleaching effects were observed.
Although not yielding II–VI cluster crystals, yet another series of crystallized
nanoparticles seems to be interesting because of a possible transfer of the synthetic
approach to the II–VI semiconductor materials also. Crystals of supertetrahedra of,
e.g., [Cd4 In16 S33 (H2 O)20 (C10 H28 N4 )2:5 ] [17] were obtained in a high-temperature
synthesis involving elemental indium and sulfur together with a cadmium salt.
Mn, Co or Zn may be introduced into the clusters instead of Cd [18]. Even larger
particles are formed without the incorporation of divalent metal ions like the cluster [In34 S54 ] 6 [19]. These structures also lack an interior core atom and form
three-dimensional superlattices together with the smaller [In10 S18 ] 6 species. A
cluster (namely [In28 Cd6 S54 ] 12 ) with the same tetrahedral shape and also with a
missing central atom is also formed by the incorporation of Cd-ions [20].
In a study dealing with the particle-particle interaction in semiconductor nanocrystal assemblies, Döllefeld et al. investigated, among other structures, the crystalline superstructure of [Cd17 S4 (SCH2 CH2 OH)26 ] depicted in Figure 4-1 [21].
In Figure 4-2 the UV-Vis absorption spectrum of this compound in solution is
shown together with the reflection spectrum of the crystalline material. It is clearly
seen that in the spectra of the crystalline material the first electronic transition
Fig. 4-2 Absorption spectrum of the [Cd17 S4 (SCH2 CH2 OH)26 ]
cluster compound in solution (dotted line) and diffuse
reflectance spectrum of the crystalline superstructure of the
same compound (cf. also Figure 4-1) [21].
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4 Organization of Nanoparticles
is shifted to lower energies by about 150 meV. Additionally, the transition is
broadened from a full width at half maximum of about 390 meV to about 520 meV.
Most probably, a complete description of the interaction of semiconductor nanocrystals in crystalline superstructures includes both electronic and dipole-dipole
interaction. The electronic coupling may be introduced by the covalent bridging of
the nanocrystals in the crystalline superstructure, but detailed theoretical studies
would help for a more quantitative picture. First attempts in this direction were
reported recently [22].
4.1.2
Superlattices of Semiconductor Nanoparticles
Yet another development of remarkable nanostructured materials yields superlattices of nanosized objects. Since there is no clear distinction between ‘‘molecular
crystals’’ and ‘‘superlattices’’ formed from nanoparticles, here we decide to refer to
structures composed of very similar but most probably not exactly identical nanoparticles, namely colloidal particles in the size range 2–10 nm.
Two excellent reviews by leading experts in the field were published in 1998 and
2000 [23, 24]. The titles of these articles contain the terms ‘‘nanocrystal superlattices’’ and ‘‘close-packed nanocrystal assemblies’’, which are in line with the aboveoutlined delimitation, although Collier et al. also report on molecular crystals as is
done here in the preceding chapter. The two articles comprise roughly 100 pages
with some 300 references and summarize the state of the art at that time in exemplary fashion. The topics include preparative aspects of the formation of monodisperse nanoparticles of various compositions including metals, the superlattice
formation itself with some theoretical background, covalent linking of nanocrystals
(see below), and an appropriate description of the physical properties and the
characterization of the nanocrystal superlattices.
Naturally, both survey papers refer to the pioneering work of Murray et al. from
1995, in which the self-organization of CdSe nanocrystals into three-dimensional
colloidal crystals is demonstrated [25] (see also [26, 27]). A prerequisite for superlattice formation is the high degree of monodispersity of the particles, a result of
synthetic advances. The characterization of the evolving structures is performed by
absorption and emission spectroscopy, high-resolution scanning and transmission
electron microscopy, electron diffraction, and, very impressively, X-ray diffraction,
giving clear-cut evidence of both long range order and tunability of the interdot
spacing by variation of the size and nature of the capping agents.
A very fine example of superlattice formation is reported by Talapin et al. [28,
29]. In Figure 4-3, the concept of the nucleation procedure is depicted: a nonsolvent (methanol) for the CdSe particles dispersed in toluene diffuses either directly
(left tube in Figure 4-3a) or slowed down through a buffer layer (propan-2-ol, right
tube) into the colloidal solution. Over weeks, this leads to irregularly shaped crystals in the case of fast diffusion (Figure 4-3b) or to perfectly faceted hexagonal colloidal crystals with sizes of about 100 mm (Figure 4-3c) consisting of about 10 12
individual 3.5 nm CdSe particles.
4.1 Semiconductor Nanoparticles
Schematic illustration of superlattice formation from
colloidal solutions of 3.5 nm CdSe nanocrystals in toluene (a).
Optical micrographs of the evolving nanocrystal superlattices
after fast (b) and slow (c) nucleation [28].
Fig. 4-3
Transmission electron microscopic images of
superlattices formed from CdSe nanocrystals. Different
orientations of the CdSe nanocrystal superlattice with
corresponding Fast Fourier Transforms as insets leading to the
assignment of fcc-like packing of the particles [28].
Fig. 4-4
In Figure 4-4, transmission electron microscopic images are displayed. Image (a)
reveals a close packing of up to seven monolayers of CdSe particles, demonstrating
the earliest stages of 3D superlattice formation starting from a 2D lattice. Images
(b) and (c) with their corresponding Fast Fourier Transforms (insets) display differently oriented segments of the evolving superlattices, leading to the assignment
of the 3D alignment of the particles to fcc-like.
There exist other approaches to the formation of superlattices aside from pure
self-organization, e.g., the use of bacterial S-layers as templates for the nucleation
of ordered two-dimensional arrays of CdS nanocrystals [30], crystallization on other
periodically patterned surfaces [31], or the ‘‘lost-wax approach’’ reported recently,
in which first polymer templates are formed from silica colloidal crystals with
subsequent filling of the evolving voids of the porous polymer with a variety of
colloidal particles [32]. Filling of voids of ‘‘artifical opals’’ (periodic structures
formed from, e.g., SiO2 , latex, or polymethylmethacrylate spheres) with semiconductor nanocrystals like CdSe [33] or CdTe [34] followed by the destruction
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4 Organization of Nanoparticles
of the template leads to so-called inverse opals, which are of eminent interest in
the evolving field of photonics. Another procedure yields CdS nanoparticles inside
3D hexagonal mesoporous silica films [35]. An area still to be approached with
II–VI semiconductor materials is the ‘‘nanotectonic’’ approach, in which threedimensional objects are formed by the template-directed assembly using nanocrystals as building units [36].
Based on work on the formation of amorphous three-dimensional arrays of CdSe
nanocrystals [37] and of compact round disks of similar particles [38], a step forward was made by Ge and Brus, who reported planar monolayer aggregates consisting of 80–100 identical 4 nm CdSe nanocrystals diffusing on planar graphite
surfaces [39]. It is, thus, an investigation into the collective motion of a very small
nanocrystal superlattice, which opens up yet another new field of research.
4.1.3
Layers of Semiconductor Nanocrystals
Layers of semiconductor nanoparticles are made in various ways, either from preformed colloids with the aid of techniques like spin-coating or layer-by-layer assembly or from ionic, atomic or molecular precursors via, e.g., electrodeposition,
chemical vapor deposition (CVD), or metal-organic chemical vapor deposition
(MOCVD). From the wealth of literature, selected examples of layered structures
are shown here as examples, together with some attempts at characterization.
Relatively early in the development of the field, the Langmuir-Blodgett (LB)
technique was recognized as a useful tool for preparing thin layers of nanocrystalline material and for controlling parameters like film composition and thickness.
Based on earlier work [40–44] in 1992, Grieser et al. utilized the LB technique for
the generation of multilayer films of CdS, CdSe and CdTe and of the mixed crystal
compounds CdSx CdSe1x , CdSx CdTe1x and CdSex CdTe1x [45]. Most of the films
are formed from multilayers of cadmium arachidate that are exposed to H2 X
(X ¼ S, Se, Te), resulting in the formation of the respective CdX particles within
the matrix of the LB films. The best characterization concerning the particle sizes
is based on the relationship between absorption energy and size, whilst direct
methods like atomic-force microscopy (AFM) and TEM do not yield unambiguous results. A fine example of an early application of a local probe technique on
colloido-chemically derived nanostructures was presented by Facci and coworkers
on similarly derived LB films of CdS particles [46]. Again, cadmium arachidate
films of various thicknesses were formed, followed by exposure to an H2 S atmosphere. The film formation was monitored by X-ray diffraction, absorption spectroscopy, nanogravimetry, and scanning tunneling microscopy (STM). In Figure
4-5 STM images from this study are shown. In the upper panel an image of a
48:5 64 nm 2 cutout of a CdS film formed from 20 bilayers of cadmium arachidate is presented, showing a structure formed by interconnected particles 4–5 nm
in size. For washing purposes the layer was treated with chloroform, which induces particle growth and aggregation, observed in the STM experiment and consistent with the absorption spectra data. From a flat zone in the upper panel a
4.1 Semiconductor Nanoparticles
STM images of a CdS film formed from 20 bilayers of
cadmium arachidate (top), small cutout of the upper panel
(middle, cf. text), and highly oriented pyrolytic graphite acting
as the substrate (bottom) [46].
Fig. 4-5
2:1 2:5 nm 2 area is displayed (middle) showing the CdS hexagonal lattice with a
period of 0.42 nm. For comparison in the lower panel, a high-resolution image of
a 2:0 1:9 nm 2 area of the supporting highly oriented pyrolytic graphite lattice is
shown, which exhibits a lattice constant of 0.246 nm. The organometallically derived TOP/TOPO-capped CdSe particles described in Section 3.1.1.3 are also suit-
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4 Organization of Nanoparticles
able to be applied as monolayer-forming elements on the water surface of an LB
trough, as demonstrated by Bawendi and coworkers [47]. The absorption and
luminescence studies indicate that the monolayers retain the optical properties of
the isolated crystallites, while TEM investigations reveal the formation of twodimensional hexagonal close-packed regions. Fendler et al. also reported on studies
of LB-derived films of CdS nanoparticles [48] and other sulfides and selenides [49].
The characterization of the films included surface pressure vs surface area isotherms as well as Brewster angle microscopy, while the focus of [48] is mainly on
the emission properties (spectral and time-resolved) of the films and that of [49] on
methodic issues of the in situ generation of size-quantized semiconductor particulate films. The emission properties of LB-derived nanocrystal monolayers are also
the field of investigations reported by Cordero et al. [50].
By comparing wet and dry films of CdSe particles with and without illumination,
it is concluded that water molecules adsorbed on the surface of the quantum dots
act firstly to passivate surface traps, accompanied by increased emission, and secondly to oxidize the surface species, introducing new surface defects which results
in a decrease of luminescence quantum yield. As demonstrated by Heath et al., the
application of the LB technique in conjunction with semiconductor nanoparticles
may lead to the generation of tunnel diodes [51]. These devices consist of a monolayer of 3.8 nm CdSe nanocrystals and an insulating bilayer of eicosanoic acid
sandwiched between an Au and an Al electrode. Advanced spectroscopic techniques like attenuated low-energy photoelectron spectroscopy were also applied
to LB-derived multilayered nanostructured assemblies of different-sized CdS particles [52].
Chemical solution deposition (CD) and electrodeposition (ED), two other techniques of film formation, make use of neutral and charged species in solution.
These sorts of film making are inseparably linked to the name of Gary Hodes. A
long list of publications could be named here (e.g. the review [53]), but instead we
refer to a few examples showing the basic ideas. The basis of CD is a reaction between a slowly produced anion (e.g. sulfide from thiourea or selenide from selenourea or selenosulfate) with a free metal cation in low concentration, which is
accomplished by complexation [54]. The nanoparticles are formed either by an ionby-ion reaction at the substrate or by homogenous nucleation in solution followed
by diffusion of the semiconductor crystallites to the substrate. Figure 4-6 displays a
number of transmission spectra of CdSe films chemically deposited at various
temperatures, with the dashed line representing particles formed at the lowest deposition temperature (274 K) and the long-short-long dashed line those formed at
the highest temperature (334 K). Obviously, smaller particles are formed at lower
temperatures, which might be explained by a quite complex reaction mechanism
involving the formation of the reacting species in solution, nucleation and growth
of crystals and the redissolution of small ones, and aggregation or coalescence of
the particles [54]. While most of the CD preparations are performed in aqueous
solution, the ED method often relies on nonaqueous electrolyes like dimethyl
sulfoxide or diethylene glycol, the chalcogen being present in its elemental form
[55, 56]. The crystallite sizes typically are in a range around 5 nm for CdS and
4.1 Semiconductor Nanoparticles
Transmission spectra of chemically deposited CdSe
films. Deposition temperatures vary from 274 K (dashed line)
to 334 K (long-short-long dashed line) [54].
Fig. 4-6
CdSe, and the properties of the films are qualitatively similar to those of the respective CD films. Differences in crystal structure are observed: whereas ED shows
hexagonal CdSe and CdS, CD shows cubic CdSe [54].
A method of film formation gaining considerable interest in recent years is
named ‘‘Layer-by-Layer’’ assembly, abbreviated LbL. The method stems from G.
Decher [57], is based on alternating adsorption of oppositely charged species, and
was originally developed for positively and negatively charged polyelectrolytes. It is
also a useful tool for the assembly of nanoparticles as shown, e.g., in a study of the
electroluminescence of negatively charged CdSe particles embedded in poly( pphenylene vinylene) [58], in work on graded semiconductor films from nanocrystals [59], and on films of IR-emitting quantum dots of HgTe [60, 61]. Because
of its universality, simplicity, and low cost, the LbL method has gained such widespread application that even including the literature on the subject would make
this chapter excessively long. Instead, only a single example is shown from the recent literature. A paper of Gao et al. [62] is devoted to lateral patterning of CdTe
nanocrystal films by a modification of the conventional LbL technique, namely by
using an electric field. The idea is that the electric field will direct the spatially selective deposition of the LbL films. Binary patterned arrays of such films containing two differently sized CdTe nanocrystals and a polyelectrolyte are obtained on
conductive indium-tin-oxide electrodes in a multistep process. Figure 4-7 displays
the impressive results of these operations. Samples a and b represent the differently sized negatively charged CdTe nanocrystals which are deposited on either the
upper (red emitting particles) or the lower (green emitting particles) biased electrode. Both sorts of particles are also found in the unbiased marginal regions,
which finally leads to orange emission in the case of successive deposition of both
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4 Organization of Nanoparticles
Photographs of lateral structures selectively formed by
green- and red-luminescing CdTe nanocrystals via electric fielddirected layer-by-layer assembly and the corresponding
emission spectra [62].
Fig. 4-7
sorts (cf. panel c and the corresponding emission spectrum on the left-hand side of
Figure 4-7).
Several other techniques have also been utilized for film formation: drop casting
(e.g., CdSe on a Pt electrode in order to study electrochromic effects in nanocrystal
films [63]), spin casting (e.g., composite films of CdSe and electron transporting
block copolymers [64]), and photochemical deposition (e.g., CdSe films from
aqueous solution [65]) and electrospray organometallic chemical vapor deposition
for the formation of quantum dot composites [66].
Regardless of the film formation procedure, the question of the interaction of
semiconductor nanocrystal assemblies arises and has been tackled by a number of
groups in recent years. In 1994, Vossmeyer et al. reported on the absorption properties of very small CdS nanoparticles [8]. A red shift of the first electronic transition with respect to the transition observed in solution is observed when compact
layers are formed from the material by spin coating. A comparatively small red
shift in the emission spectra of the crystalline superstructure of CdSe nanoparticles, mentioned in Section 4.2.2, with respect to the emission peaks in solution is explained by small size deviations of the particles within the sample result-
4.1 Semiconductor Nanoparticles
ing in energy transfer from the smaller to the larger particles [25, 67, 68]. Similar
observations were made for III–V semiconductor nanocrystals of InP [69]. Artemyev et al. prepared unstructured solids also from CdSe nanoparticles with polymers acting as spacers in order to control the distance between the particles
[70–72]. The pure solid built up by drop casting solutions of semiconductor nanoparticles shows very broad transitions bearing a resemblance to a bulk spectrum
rather than to isolated absorption bands. Micic et al. obtained similar results in
disordered films of InP nanoparticles [73]. After studies on the photoconductivity
in CdSe quantum dot solids [74], Leatherdale et al. reported on CdSe particles
in different dielectrics [75]. In agreement with the results in [8], a red shift of
the first electronic transition was found exhibiting no significant broadening. Kim
et al. presented studies of pressure dependency in solutions and compact layers of
semiconductor nanoparticles of CdSe [76]. Different stabilizing agents were used,
and, in contrast to the results of Leatherdale et al., no red shift was found for the
compact layers of TOP/TOPO-stabilized CdSe particles. Only for particles stabilized with pyridine was a shift to lower transition energies observed. Again, no
significant broadening of the electronic transitions was noted. As shown in Section 4.2.1, Figure 4-2, reflection spectroscopy on the cluster crystal compound
Cd17 S4 (SCH2 CH2 OH)26 reveals substantial broadening and a low-energy shift of
the first absorption band, which in a first step is explained by both dipole-dipole
interaction and electronic coupling of the nanocrystals in the crystalline superstructure. When comparing these results with those from disordered solids obtained by spin-coating concentrated solutions of the same material, it was found
that extreme care has to be taken in order to avoid measuring artifacts [77]. Misinterpretations of observed alterations of transition energies in these structures are
avoided by the choice of an integrating sphere as the experimental setup. The resulting ‘‘real’’ red shift of the transition energy is explainable in a simple picture by
dipole-dipole interactions of the semiconductor nanocrystals in the films. More
elaborate theoretical work covers topics like exciton dissociation and interdot
transport [78], takes into account the dielectric medium surrounding the nanoparticles [79], and explains an observed solvatochromism in CdSe colloidal quantum dots, also taking into account the screening of the ligand shell [75], and, for
example, uses pair potentials to predict the phase behavior and thermodynamics of
small particles [80]. Because of its potential applications (based also on experimental studies of, e.g., fast kinetics [81] and electronic transport in films of colloidal particles [82–86], this research area will remain active in the coming years.
4.1.4
Coupling of Semiconductor Nanocrystals
Probably one of the most demanding challenges in current preparational research
is the directed coupling of semiconductor nanoparticles. If the quantum dots were
called ‘‘artificial atoms’’, this endeavor would yield ‘‘artificial molecules’’, and the
challenge is to create such quantum dot molecules in a desired manner to form,
e.g., dimers (‘‘two-atomic quantum dot molecules’’) of various compositions, AB4
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4 Organization of Nanoparticles
molecules, etc. The field is about to grow out of its infant status, and so the very
first attempts in this direction are reported here.
Couplings of semiconductor nanocrystals to, e.g., small particles of metals [87]
or insulators [88, 89] is beyond the scope of this chapter. Additionally, couplings
of semiconductor nanoparticles to DNA [90–94] and to other biological systems
[95–100], as well as those to organic [101] or metal-organic fluorophores [102], are
excluded here, and the literature can therefore be more clearly arranged.
A recent paper by Kolny, Kornowski and Weller reports on the use of the electrostatic interaction between oppositely charged CdS nanocrystals [103]. It is
shown by absorption spectroscopy, X-ray diffraction, and high-resolution electron
microscopy that two types of aggregates may be formed: complex-like structures
are obtained when a large excess of one kind of particles is used, whereas a threedimensionally ordered solid evolves when comparable amounts of the constituents
are mixed. Additionally, the strength of the interaction between the particles is
controllable by adjusting the ionic strength of the solution. Another approach relies on the covalent bonding of different [104, 105] or similar [106] semiconductor
nanocrystals. The latter work describes the synthesis, isolation and characterization of CdSe-homodimers. Monodisperse CdSe particles capped with TOPO were
treated with various surface-active species. Thiols proved to be a good choice for
passivation of the surfaces. For coupling of the nanocrystals, bis(acylhydrazine)
was used, a compound bridging two semiconductor nanocrystals via its two terminating NH2 -groups. Size-selective precipitation finally leaves mostly homodimers
in solution. The covalent coupling of different CdTe nanocrystals is reported in
[104]. The CdTe particles are stabilized with thioacids or aminothiols, which form
an amide bond using carbodiimide as a mediating agent. Thus, in this case, the
stabilizing ligands are directly used for the coupling without a bridging molecule,
in contrast to the above-mentioned case and the work of Trindade et al. [105], in
which the coupling of PbS particles to CdS or CdSe nanocrystals via the bridging
ligand 2,2 0 -bipyrimidine is outlined. Figure 4-8 displays the infrared transmission
spectra of a sample containing only a mixture of the two kinds of nanocrystals
(dotted line) and of a sample of the coupled particles (solid line). A careful analysis
of these spectra, together with those of the bare particles, reveals the existence of
the NaH vibrations within the amide group and the diminishing of the OaH
vibration in the carbonyl group due to the coupling reaction (see the arrows in
Figure 4-8). The authors also state that a possible alternative interpretation needs
to be considered. In their own words: ‘‘Since the nanocrystalline powders are
washed with buffer solution and dried in a vacuum in order to perform the IR
measurements, the existence of coupled stabilizers without connection to at least
one nanocrystalline surface can be excluded. However, it is conceivable that some
of these ‘‘amide bridges’’ fall off the nanocrystal on one side. Since the SaH bond
vibrations are detected in neither Raman nor IR spectroscopy, this hypothesis cannot yet be verified.’’
Together with theoretical investigations [22, 107, 108], this sub-field of nanoscience is expected to continue and to grow as more sophisticated preparational
attempts are launched and new demands from the applications area are imposed.
4.2 Metal Nanoparticles
Fig. 4-8
Infrared transmission spectra of mixed and linked CdTe nanocrystals [104].
4.2
Metal Nanoparticles
Günter Schmid, Dmitri V. Talapin, and Elena V. Shevchenko
The organization of metal nanoparticles is, like that of semiconductor nanoparticles, of vital relevance. The ultimate goal of all efforts in making and investigating
semiconducting or metal nanoparticles is their use in optoelectronics, nanoelectronics, storage systems, etc., based on the size-quantized properties of these miniaturized materials. The fact that the electronic properties of metal particles dramatically change if their size falls below a well-defined limit makes them very
promising candidates for single-electron switches, transistors, etc. Organization of
the particles in three (3D), two (2D) or even in one (1D) dimension(s) seems indispensable if specific functions are be called upon. This chapter therefore reports
on the present state of the art with respect to 3D, 2D, and 1D arrangements of
metal nanoparticles.
4.2.1
Three-dimensional Organization of Metal Nanoparticles
In principle, 3D organization of building blocks of matter happens simply by
nature: we call it crystallization. Why do we have to devote a special chapter to 3D
organization of metal nanoparticles? Typically, ions and molecules forming 3D
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crystals are identical in every respect and so enable them to be arranged periodically in three dimensions. This is usually not the case for building blocks on the
nanometer scale. Nanoparticles, traditionally also called ‘‘colloids’’, usually differ
not only in size, but also in shape and structure. The surprising discovery with respect to crystal formation of nanoparticles during the last decade was that, in spite
of not having fully identical building blocks, 3D organizations and (as will be
shown in the following sections) 2D and even 1D arrangements become available
if specific conditions are fulfilled. Of course, there is a continuous transition from
1D to 2D and from 2D to 3D. In the following, we will consider assemblies of
some monolayers as quasi-two-dimensional systems or thin films rather than as
3D architectures. Three-dimensional crystals are considered to exhibit dimensions
in the a, b, and c directions of the same order of magnitude.
Usually, nanoparticles to be 3D assembled to superlattices are covered by a shell
of stabilizing ligand molecules, protecting the particles from coalescence and, furthermore, making them soluble in appropriate solvents.
Gold nanoparticles, stabilized by alkylthiolates RS (R ¼ n-C12 H25 ) and rather
narrowly distributed in size, could be obtained by repeated fractional crystallization
of a mixture of 1.5–3.5 nm particles from nonpolar solvents. The success of the
fractionation was controlled by mass spectroscopy [109]. Based on to these data,
definite numbers of atoms could be assigned to the particles of distinct fractions,
e.g., 140, 225, 314, and 459, all consisting of fcc-structured octahedra. Solutions of
these remarkably well-defined particles can be used for generating 2D assemblies
(see Section 4.2.2) and also for growing 3D crystals of considerable extension. The
Au459 fraction forms well-shaped crystals up to ca. 40 mm in length [109]. Obviously no coalescence is happening during crystal formation, since the crystals are
highly soluble in nonpolar solvents. The assembly of nanoparticles changes from
the well-known purple color of individual colloids to a metallic-like luster.
Superlattices of 5.0 nm silver nanoparticles have also become available from
alkylthiolate protected colloids [110–112]. From high-resolution electron microscopy (TEM), quite precise information about the architecture of the crystals could
be deduced. They consist of highly ordered truncated octahedra, as is formally
shown in Figure 4-9.
Crystals of up to ca. 5 mm edge length could be grown on the TEM copper grid
from solution. 3D superlattices of thiol-stabilized 5–6 nm gold particles have also
been described [113].
Micron-sized 3D crystals of the well-known gold cluster compound Au55 (PPh3 )12 Cl6 are formed from dichloromethane solution, simply by evaporating the
solvent [114]. This is no surprise, since these very small nanoparticles are indeed
identical in size and shape. The growth of larger crystals fails because of the fast
decomposition of these clusters in solution in the course of hours. The fast removal of the solvent results in small crystals, which, however, are more or less
perfectly organized, as can be followed from TEM and small-angle X-ray diffraction
(SARXD) measurements. The 2y angle of 4.3 corresponds with a d value of 2.05
nm, in agreement with an fcc arrangement of the clusters of an effective size of
2.3–2.4 nm including the ligand shell. From TEM, the hexagonally close-packed
4.2 Metal Nanoparticles
3-D
{100}
<011>
{ 111
}
{ 1 11 }
{100}
{ 11 {100}
}
1}
11
{1
[ 110] s
[ 110]s
Fig. 4-9
Sketch of a superlattice of truncated octahedral 5.0 nm Ag nanoparticles.
arrangement of the clusters can also be deduced, in agreement with the SAXRD
result.
Independent of the dimensionality, the organization of metal nanoparticles
seemed to be linked to the existence of a protecting ligand shell enveloping the
particles’ surfaces. The most important reason for this condition was seen as the
necessity to grow the arrangements from solution.
The soft removal of the ligands from protected particles in solution could be one
way to realize the generation of defined naked nanoparticles. On the other hand, it
is known that a quantitative elimination of a ligand shell cannot be induced by
heat treatment. As early as 1986, the first hints emerged that naked metal clusters
could organize themselves in an ordered manner. Mild electrochemical decomposition of Au55 (PPh3 )12 Cl6 in solution resulted in the formation of naked Au13 clusters (the inner cores of Au55 ) together with Ph3 PAuCl and bulk gold. The Au13
clusters were found to build up clusters of clusters of the type (Au13 )n [115, 116].
Secondary ion mass spectrometric (SIMS) investigations of Au55 (PPh3 )12 Cl6 resulted in similar species, identified by masses corresponding to (Au13 )n with n
values up to 55 [117, 118].
The generation of bare Au55 clusters under mild conditions has later become
possible by a surprising reaction: a thiol-terminated fourth generation dendrimer
with 96 SH groups (G4-SH) on its surface reacts with Au55 (PPh3 )12 Cl6 in solution
with quantitative elimination of PPh3 and Cl ligands [119].
The reason for the removal of the ligand shell is to be seen in a formal substitution of the phosphines by the stronger thiol groups of the dendrimers, which are
used in large excess. Ligand-free clusters on the dendrimers’ surface are obviously
not irreversibly fixed, but can move on the equivalent surface positions to finally
get contact with a second, third etc. other bare cluster, with formation of strong
metal-metal bonds. The process can be described as a crystal growth in a matrix of
253
254
4 Organization of Nanoparticles
Fig. 4-10 Illustration of the formation of Au55 microcrystals
inside a matrix of G4-SH dendrimers.
Fig. 4-11
TEM image of an Au55 microcrystal with a skin of G4-SH dendrimer molecules.
organic material. The growth continues until the solubility is exceeded. Figure 4-10
illustrates in a very simplified manner the possible formation of the isolated microcrystals consisting of only Au55 building blocks. One of these perfectly shaped
crystals is shown in Figure 4-11.
Surprisingly, not only are the PPh3 ligands removed and can be detected in solution, but also the Cl atoms have disappeared, recognizable in the lack of the
characteristic AuaCl vibration at 274 cm1 in the IR spectrum.
The most important question is, how do the bare Au55 clusters organize themselves to give the perfect-looking microcrystals? Detailed TEM and small- and wideangle X-ray diffraction (SAXRD, WAXRD) finally showed that the cuboctahedrally
shaped Au55 clusters obviously touch each other via their edges in a very regular
4.2 Metal Nanoparticles
Fig. 4-12 Magnified TEM image of an Au55 microcrystal and
crystallographic data from SAXRD and WAXRD. The hkl values
111 and 110 are reduced by 5% because of size-determined
contraction of AuaAu distances in small particles.
manner to build up superlattices of an architecture schematically shown in Figure
4-12.
These results show impressively that the organization of metal nanoparticles
need not be limited to ligand-protected ones, but that ligand-free particles can also
be three-dimensionally organized to build up new modifications of metals, in this
case gold.
In contrast to the microcrystals formed of Au55 (PPh3 )12 Cl6 , which have been intensively investigated with respect to their electronic characteristics, the 3D assemblies of bare Au55 clusters have up to now not been subject of physical investigations.
Self-assembly of monodisperse nanoparticles into ordered 3D superstructures is
usually achieved by slowly destabilizing a colloidal dispersion. Indeed, if the evaporation rate of a solvent is relatively high, close-packed glassy films of nanoparticles
with short-range order are formed [120]. On the other hand, superlattices with low
concentrations of lattice defects and a long-range order of nanocrystal arrangement
can be formed if the particles are mobile and have enough time to find their lowest
energy sites in the superstructure. For example, slow evaporation of a high-boiling
carrier solvent at elevated temperatures (50–60 C) was used to form ordered su-
255
256
4 Organization of Nanoparticles
Fig. 4-13 Schematic outline illustrating the concept of the
three-layer oversaturation technique used for the crystallization
of monodisperse FePt and CoPt3 nanoparticles.
perstructures of monodisperse cobalt [120–123], nickel [121], and magnetite [124]
nanoparticles. Recently, the technique for growing macroscopic crystals (so-called
colloidal crystals) consisting of monodisperse semiconductor or metal nanoparticles as building blocks has been developed [125–127]. The method is based on
gentle destabilization of a colloidal solution induced by slow diffusion of a nonsolvent into a concentrated solution of monodisperse nanoparticles inside a vertically positioned glass tube, as shown in Figure 4-13. The colloidal crystals nucleate
and grow preferentially on the walls of a glass tube. The growing process usually
takes several weeks. The spatial distribution of local oversaturations caused by the
diffusion of a nonsolvent determines the quality of the colloidal crystals. To make
the oversaturation front less sharp than in the case of a direct solvent-nonsolvent
contact, a buffer layer with low, but not negligible, solubility for nanocrystals was
inserted between solvent and nonsolvent layers. This technique allowed preparation of perfectly faceted colloidal crystals with sizes from tens of microns up to a
hundred microns from monodisperse CdSe [125], FePt [126], and CoPt3 [128]
nanocrystals.
Figures 4-14a and b show optical micrographs of FePt colloidal crystals adhering
to the glass. The crystals grew preferably in the form of faceted triangular or hexagonal platelets, 10–30 mm in size, on a slide. Crystals with a pyramidal-like shape
were also observed. Scanning electron microscopy (SEM) provides a closer look at
the morphology of colloidal crystals. In the case of FePt, faceted triangular platelets, incomplete hexagonal platelets, and tetrahedral colloidal crystals with imperfect sides showing terraces and cleaved ledges (Figure 4-14c–e) were clearly distinguishable. Figure 4-14 shows a TEM image taken from a small fragment of the
4.2 Metal Nanoparticles
Fig. 4-14 Optical micrographs (a, b) and SEM images (c, d, e)
of colloidal crystals of FePt nanocrystals illustrating their
different shapes. TEM image (f ) of the fragment of an FePt
colloidal crystal. Reproduced from [126], Copyright 2001, with
permission from Wiley-VCH.
257
258
4 Organization of Nanoparticles
SEM images of colloidal crystals of CoPt3
nanoparticles illustrating their different shapes. Reproduced
from [128], Copyright 2002, with permission from the American
Chemical Society.
Fig. 4-15
Fig. 4-16
HRSEM images of the surface of a CoPt3 colloidal crystal.
FePt colloidal crystal obtained by mechanical grinding and treatment in an ultrasonic bath. Hexagonal arrangements of nanocrystals, i.e. building blocks of the
colloidal crystals, are clearly seen at the edges of the crystalline pieces. Figure 4-15
shows SEM images of well-faceted hexagonal, triangular, and pyramid-like colloidal
crystals built of CoPt3 nanocrystals.
The arrangement of nanocrystals at the surface of colloidal crystals can be investigated using HRSEM (Figure 4-16). Thus, the ledges and terraces of ordered
CoPt3 nanoparticles are clearly seen at the surface of colloidal crystal. These 6.3
nm CoPt3 nanoparticles are superparamagnetic, with the blocking temperature of
19 K. However, in 3D ordered structures dipolar interactions between adjacent
4.2 Metal Nanoparticles
particles can orient the magnetic moments of individual nanocrystals, and ferromagnetic domains can appear within the supercrystal resulting in considerable
increase of the blocking temperature [123, 129].
The self-assembly of magnetic nanocrystals can also be affected by applying an
external magnetic film. We have found that relatively large external magnetic fields
disturb the crystallization of magnetic nanoparticles into macroscopic colloidal
crystals, causing the formation of glass-like solids rather than 3D ordered superlattices.
The direction of the applied magnetic field influences the morphology of selfassembled structures of magnetic nanoparticles [121, 130, 131]. When the field
was applied in the plane of the substrate, deposition of Co nanoparticles led to
spindle-shaped superlattices with their long axes aligned along the external magnetic field [121, 132]. Within these spindle-shaped aggregates, individual magnetic
nanoparticles showed nearly perfect long-range order [121]. CoPt3 nanocrystals
slowly precipitated in the external magnetic field of 0.5 T applied parallel to the
substrate, organizing themselves into parallel equidistant stripes, as shown in Figure 4-17a. In such structures, the periodic stripes of close-packed CoPt3 nanocrystals were separated from each other by regions where no nanocrystals were
deposited. The period of these self-assembled gratings decreased with increase in
the strength of the external magnetic field.
A magnetic field applied perpendicularly to the substrate induces the formation
of 2D hexagonal superlattices of individual Co nanoparticles [132]. Pileni et al.
[130] observed the formation of a hexagonal network of ca. 1 mm dot-shaped aggregates made of 8 nm cobalt nanoparticles. CoPt3 nanocrystals (4 nm) deposited
under the magnetic field of 0.9 T applied perpendicularly to the substrate can selfassemble into needles of diameter 0.1–2 mm and length up to 100 mm (Figure
4-17b). HRSEM investigation revealed no long-range ordering of CoPt3 nanoparticles inside these needles.
Self-organization phenomena in colloids of magnetic nanoparticles (ferrofluids)
induced by external magnetic fields have attracted considerable attention because
of their importance from both fundamental and an applied perspectives [133].
However, the lack of systematic investigations and some contradictory published
results can be noticed, probably because of the relative novelty of this subject.
Wirtz et al. [134] suggested that the formation of macroscopic one-dimensional
periodic patterns composed of high- and low-concentration regions in a ferrofluid
is possible only in an oscillating magnetic field. On the other hand, Wang et al.
succeeded in forming periodic branched structures similar to those shown in
Figure 17a under a constant magnetic field [135]. Recently, Wen et al. modeled the
self-assembly of planar magnetic colloidal crystals under an external magnetic field
using glass microspheres coated with nickel floating at the interface of two immiscible liquids of different densities [136]. A hexagonal lattice with the lattice
constant steeply increasing with increase of magnetic field strength was observed
when the magnetic field was applied perpendicularly to the plane of the magnetic
particles. Tilting the magnetic field at an angle of 60 away from the normal projection transformed the hexagonally ordered monolayer into equidistant straight
259
260
4 Organization of Nanoparticles
Fig. 4-17 a Stripes of 9 nm CoPt3 nanocrystals deposited in a
magnetic field of 0.5 T applied parallel to a glass substrate. b
Needles composed of 4.0 nm CoPt3 nanocrystals deposited in
the magnetic field of 0.9 T applied perpendicularly to a glass
substrate. The image was acquired at a 40 tilt of the substrate.
chains of particles aligned along the field [136]. Studying self-organization phenomena in ferrofluids requires extremely monodisperse samples of magnetic
nanoparticles, because the magnetic dipole moment of a single-domain nanoparticle is proportional to the third power of the radius, and therefore particles of different sizes have different susceptibilities to the magnetic field. The considerable
progress in the preparation of monodisperse magnetic nanocrystals achieved during recent years will probably help our further understanding self-organization and
phase behavior of ferrofluids.
4.2.2
Ordered Monolayers of Metal Nanoparticles
Nature tends to organize materials three-dimensionally. This is why 2D assemblies
usually do not grow except within special experimental conditions. As already
mentioned in the previous section on 3D structures, it turned out that formation of
4.2 Metal Nanoparticles
2D assemblies can be achieved without having identical building blocks. More than
that, the tolerance with respect to deviations from size and shape is found to be
even larger than in the 3D case. The formation of 2D structures of metal nanoparticles has therefore developed into a broad field in nanoscience, with promising
aspects regarding future applications.
The present section deals with the different techniques of making ordered
monolayers. Double and multiple layers up to thin films will also be considered
here, since these systems are not covered by the definition of 3D crystals given in
Section 4.2.1.
The generation of ordered 2D structures of metal nanoparticles can be achieved
by very different techniques:
– self-assembly from solution
– use of pre-prepared templates
– writing structures with special techniques.
In the following, examples of these three principal routes will be presented and
discussed.
Self-Assembly
Self-assembly of nanoparticles on appropriate substrates is, without doubt, the
simplest way to make monolayers. However, there are several disadvantages associated with this method: there is little or no influence on the structure and extension of the monolayers. Formation of self-assembled monolayers (SAMs) depends
on experimental conditions, making reproducibility difficult. Nevertheless, many
quite impressive 2D structures of metal nanoparticles have become known in the
course of the last decade.
One of the first hints that gold nanoparticles can be deposited in an ordered
manner came from Schiffrin et al. [137]. Gold particles (8 nm), prepared in a twophase liquid-liquid system and stabilized by dithiols [138], gave cross-linked assemblies whose quality of ordering was rather limited compared with recent
results. However, these very first results showed that 2D self-assembled arrays can
be produced.
This research group and others contributed considerably to the improvement of
2D architectures by varying the experimental conditions [139]. Instead of dithiols,
tetraalkylammonium bromides [140], [R 4 N]þ Br, are used, whereby the separation
of the clusters depends on the chain length of R in [R 4 N]þ . Bilayered structures
also became available with these particles, forming unusual patterns owing to the
occupancy of 2 gold saddle sites, and leading to linear and circular arrangements.
The reason for this unexpected structure is to be seen in the balance between local
electrostatic repulsion and dispersion forces between the ammonium stabilized
particles.
Schiffrin et al., also for the first time, made another important discovery:
they fabricated bimodal ensembles of nanosized gold particles. Alkanethiolstabilized gold nanoparticles of different but well-defined sizes can organize two4.2.2.1
261
262
4 Organization of Nanoparticles
dimensionally, namely with size segregation, each region consisting of hexagonally
close-packed structures, randomly positioned particles, and, most importantly, ordered bimodal arrays also.
‘‘Alloy-like’’ bimodal 2D structures, consisting of two different metals, are likewise available, as was shown for silver and gold [141, 142]. There seems to be a
relationship between the geometrical packing of these sphere-like particles and
that of intermetallic alloys on an atomic scale. This follows the well-known Laves
rules [143], which describe the ratio of the atomic radii rA and rB of two metallic
species rA =rB (where rA > rB ) and of the local number ratio of the two metals
nA =nB . Furthermore, there are polarizability components to be considered. In the
case of only one metal of two different particle sizes (< 10 nm), the 2D lattice
structures mainly depend on the ratios nA =nB and rA =rB , since polarization effects
can be neglected. AB2 structures are formed if nA =nB ¼ 0:5 and 0:482 < rA =rB <
0:624. If, however, nA =nB ¼ 1 and 0:27 < rA =rB < 0:425, AB lattices with the NaCl
structure result. Both alternatives are sketched in Figure 4-18.
An Au/Ag mixture whose particles are of approximately the same size gives only
randomly oriented colloids. However, an Au/Ag 10:1 mixture with 8 nm Au and 5
nm Ag particles results in an AB2 superlattice, in agreement with the Laves alloy
rules. It should be stated that these comparisons of true alloys with mixtures of
nanoparticles of different sizes are not really permissible, since the conditions do
not correspond. For instance, the fact that the nanoparticles are ligand-stabilized
means that the metal cores do not really touch each other. However, as a simple
model to understand the different structures of the 2D superlattices it works quite
satisfactorily.
When two monodisperse colloids of CoPt3 nanocrystals (4.5 nm and 2.6 nm in
diameter) were mixed together, followed by slow evaporation of the solvent, an
AB5 -type superlattice analogous to the structure of the intermetallic compound
CaCu5 [144] was obtained (Figure 4-19) [128]. A similar structure was observed for
binary mixtures of latex spheres of two different sizes [145, 146]. In the first plane
of this lattice (Figure 4-19), each 4.5 nm CoPt3 nanocrystal is surrounded with a
hexagon formed by 2.6 nm nanocrystals. The second plane consists only of hexagons of small particles, and the third plane repeats the first.
Well-ordered and largely extended (up to 1 mm) 2D superlattices formed by selforganization of silver and gold nanoparticles protected by alkylthiolates were found
by Whetten et al. in connection with the formation of 3D assemblies (see Section
4.2.1) [109, 110]. The Ag particles were produced by a special aerosol technique in
which elementary silver was evaporated at 1200–1500 K into an atmosphere of
helium, followed by a cooling step to give 4–5 nm Ag particles. These were finally
cocondensed with an excess of alkylthiols (C12 chains) [110].
The gold nanoparticles were the same as those used for the fabrication of 3D
crystals [109]. Fcc-structured monolayers and multilayers were formed, from solution, on carbon-coated copper grids. Detailed structural investigations are described
in the original literature.
Other groups also contributed to this development in the second half of the
1990s, mostly also using the noble metals silver and gold [147–149, 113].
4.2 Metal Nanoparticles
Fig. 4-18
Sketch of 2D AB2 (a) and AB (b) superlattices.
Spontaneous self-assembly was observed when colloidal solutions of CoPt3
nanocrystals were spread onto a substrate, with subsequent slow evaporation of the
carrier solvent [127, 128]. Depending on the particle size and conditions of solvent
evaporation, several kinds of self-organized superstructures were obtained. Figure
4-20a shows a TEM overview image of a close-packed monolayer of 4.8 nm CoPt3
nanocrystals. If the surface coverage with nanocrystals was thicker than one monolayer, nanocrystals of the second layer occupied positions between the nanocrystals
in the first layer (Figure 4-20b, c). Owing to the relatively large interparticle spacing
263
264
4 Organization of Nanoparticles
Fig. 4-19 a Schematic view of an intermetallic compound
CaCu5 . b, c TEM images illustrating AB5 -type superstructure
formed by CoPt3 nanocrystals with bimodal size distribution.
Reproduced from [128], Copyright 2002, with permission from
the American Chemical Society.
(ca. 2.5 nm) maintained by the bulky ACA capping ligands, two layers of CoPt3
nanocrystals formed the above mentioned structure, observable for different nanocrystal sizes. The third layer of nanocrystals occupied the positions typical of the
cubic close packing (ccp) structure (Figure 4-20d, e). Figure 4-20e shows a magnified image of a three-layered assembly of ca. 4.0 nm CoPt3 nanocrystals. The difference between the phase contrast of the two underlying layers and that of the
darker third layer allows us to attribute each nanocrystal to the layer it is placed
in and to conclude that CoPt3 nanocrystals are packed into a ccp-like superlattice
where the nanocrystals are separated from each other by thick (2.5 nm) organic
shells. Figure 4-20f shows a TEM image of a multilayer 3D superlattice where the
4.5 nm CoPt3 nanocrystals are arranged in a nearly defect-free 3D structure, ex-
4.2 Metal Nanoparticles
Fig. 4-20 TEM images of (a) a monolayer of
4.8 nm CoPt3 nanoparticles; b and c two layers
of 3.6 and 4.0 nm CoPt3 nanoparticles; d three
layers of 3.6 nm CoPt3 nanoparticles; e three
layers of 4.0 nm CoPt3 nanoparticles; f more
than five layers of 4.5 nm CoPt3 nanoparticles;
g graphical illustration of a three-layer
arrangement of CoPt3 nanocrystals.
Reproduced from [128], Copyright 2002, with
permission from the American Chemical
Society.
hibiting long-range order. The graphical illustration (Figure 4-20g) clarifies the
three-layer arrangement of CoPt3 nanocrystals.
The shape of the individual nanocrystals can strongly affect the morphology of
self-assembled structures [150–152]. In the case of spherical particles, the crystalline lattices of individual nanocrystals are as a rule randomly oriented within a
superstructure. However, faceted and, in particular, cubic nanocrystals can form
‘‘lattice-matched superstructures’’, in which each nanocrystal has the same orientation as its neighbors (Figure 4-21). Such superstructures are of interest for creating materials with a high magnetic anisotropy constant, because of the possibility
of aligning the axes of easy magnetization of individual nanocrystals.
The high magnetic anisotropy of CoPt3 nanocrystals [153] makes them superior
candidates for realizing this new kind of artificial solid.
Another exciting example of the effect of particle shape on self-assembly is the
formation of the cobalt nanorod superstructures observed by Chaudret et al. [154].
Exceptionally monodisperse cobalt nanorods were synthesized by thermal decomposition of Co(h3 -C8 H13 )(h4 -C8 H12 ) organometallic precursor in anisole at 150 C
in the presence of hexadecylamine and stearic acid under hydrogen at a pressure of
3 bar. Upon aging the colloidal solution, the Co nanorods stuck to each other,
forming a superstructure consisting of perfectly aligned nanorods (Figure 4-22).
265
266
4 Organization of Nanoparticles
Fig. 4-21
TEM images of self-assembled structures of cubic CoPt3 nanocrystals.
Fig. 4-22 3D organization of hcp cobalt nanorods. Side and
top views. Courtesy of Dr. B. Chaudret.
Note that the axes of easy magnetization of individual nanorods are aligned in one
direction, so that one would expect high magnetic anisotropy and very high coercivity of the material obtained.
The wetting properties of the substrate can influence the superlattice morphology [121]. If the colloidal solution of nanoparticles wets the substrate, a 2D superlattice grows preferentially, forming a monolayer. As the surface coverage in-
4.2 Metal Nanoparticles
creases, the nanocrystals adsorb to the ledges and kinks of the growing structure,
forming terraces that extend laterally across the substrate. If the colloidal solution
does not wet the substrate, a 3D superlattice grows preferentially, with facets
reflecting the packing symmetry inherent in the nanocrystals. Annealing of the
superlattice can change the packing symmetry of self-assembled magnetic nanocrystals. Upon annealing of 2D arrays of e-Co nanoparticles under vacuum or inert
atmosphere at ca. 300 C, the packing symmetry of 3D ordered regions was converted from the original hexagonal to cubic [122]. The internal transition of the
nanocrystal structure from cubic e-Co to hcp-Co was responsible for this change.
Further annealing of Co nanoparticles led to dramatic change of the crystalline
core, fusing and sintering of nanoparticles, and destruction of the superlattice. The
temperature of nanocrystal coalescence may depend on the substrate. Thus no
obvious aggregation of PtFe nanocrystals in a monolayer was observed even at
700 C on an SiO2 substrate, while on an amorphous carbon surface the particles
coalesced at ca. 530 C [150, 155].
The monolayers described so far are all grown by self-assembly processes where
ligand-stabilized metal nanoparticles are deposited on ‘‘innocent’’ surfaces from
solution, and this is followed by a more or less fast evaporation of the solvent. Depending on concentration and other experimental conditions, those procedures
may sometimes also result in multilayers or even in small 3D microcrystals. Selfassembly processes can, however, also be supported by chemical, electrochemical,
magnetic, or mechanical effects. Some examples of these are given in the following
sections.
The organization of metal nanoparticles on surfaces is based not only on the
presence of almost identical building blocks and attractive forces between them,
but also on the interaction between particles and surface. Strong bonds between
nanoparticle and surface do not usually lead to ordered structures of any appreciable extension. The reason is that chemically bound particles cannot move sufficiently to support the organization process. For instance, only very small areas
of ordered Au55 (PPh2 C6 H4 SO3 H)12 Cl6 clusters could be obtained on poly(ethyleneimine)-modified surfaces, since the relatively strong acid-base interactions between clusters and surface disturb the organization process [156]. On the other
hand, the lack of any interaction between surface and particle may also be disadvantageous, since there is no reason to deposit from solution as long as solvent
is present. These conditions often end up in the formation of circles of particles
owing to the drying mechanism of the solvent. It turned out that moderate attraction between nanoparticles and surface are best suited to generate ordered areas of
remarkable extent. The following examples illustrate these findings.
Au55 (PPh3 )12 Cl6 , dissolved in dichloromethane, can be distributed in a very thin
film on a water surface. The particles self-assemble on this surface, resulting in
relatively small hexagonal or square structured islands that can be transferred to
TEM grids for imaging [157]. Figure 4-23 shows some of these small islands. This
self-assembly process can be drastically improved if the attraction between the
surface and the nanoparticles is slightly increased. A monomolecular film of appropriate polymer molecules at the water-dichloromethane phase transition sup-
267
268
4 Organization of Nanoparticles
Fig. 4-23 TEM image of Au55 (PPh3 )12 Cl6 clusters, selfassembled on a water surface from dichloromethane solution.
ports the attraction. Molecules with an extended p electron system, interacting with
the electronic p system of the phenyl rings of the PPh3 ligands, turned out to work
best. Examples are poly( p-phenyleneethynylene) (PPE) derivatives, but poly(vinylpyrrolidone) (PVP) and several other polymers gave comparable results [158, 159].
In addition, these polymer films act as a kind of ‘‘carpet’’, supporting the transfer
to solid substrates. Figure 4-24 impressively demonstrates the difference between
self-assembled areas without (Figure 4-23) and with polymers.
The relevance of the weak particle-surface interaction could also be demonstrated when alkyl- instead of aryl-substituted ligands were used: there was no
ordering process at all. However, if the same clusters with an ‘‘alkyl surface’’ were
used in combination with saturated polymers such as poly(methylmethacrylate)
(PMMA), reasonable ordering could be observed again [157].
Independently of the polymer, two different arrangements of Au55 (PPh3 )12 Cl6
were observed: hexagonal and square. The reason for this is not quite clear; however, it can be postulated that the very first contact between a few clusters will
be decisive for the formation of the one or the other structure, since the orientation of the PPh3 ligands between different clusters could determine the further
growth. Figure 4-25a illustrates the situation schematically, while Figure 4-25b
shows a simulation of a full space model, elucidating the different densities [160].
Indeed, statistics calculated from many different TEM images clearly show different cluster-cluster distances, depending on the structure of the monolayers, as indicated in Figure 4-25a.
The TEM image in Figure 4-26 shows a cutout of a perfectly square-packed
monolayer of Au55 (Ph2 PC6 H4 SO3 H)12 Cl6 clusters [158].
Guided Self-Assembly
A ‘‘guided’’ self-assembly of gold nanoparticles on patterned GaAs surfaces has
been elaborated by R. P. Andres et al. for developing arrays electronically to
4.2.2.2
4.2 Metal Nanoparticles
Fig. 4-24 TEM image of a cutout of extended self-assembled
Au55 (PPh3 )12 Cl6 arrangements, deposited on a polymer film as
support.
form semiconductor device layers [166]. A structured Be-doped GaAs layer is covered by a monolayer of xylyl dithiol, HSaCH2 aC6 H4 aCH2 aSH. If a solution of
dodecanethiol-protected 5 nm Au nanoparticles in hexane is spread on a convex
water surface, a uniform monolayer of nanoparticles is formed, in accordance with
the procedure described above. If the modified GaAs surface is caused to briefly
touch the nanoparticle monolayer, this is transferred to the substrate. It is assumed
that dodecanethiol ligand molecules are partially substituted by the xylyl dithiol
covering the surface. The quality of the hexagonal close-packed monolayer of gold
nanoparticles was controlled by transferring samples to TEM grids. It can be assumed that the quality of the monolayer does not change if the thiol-modified
GaAs surface is used instead of the grid. This procedure could be of particular interest, since the self-organization happens on a water surface, and the monolayer is
then strongly fixed onto another surface by covalent bonds. Rinsing of the surface
by hexane only removes particles that are not tethered, but not the chemically fixed
ones.
Langmuir-Blodgett (LB) techniques have also been used to fabricate ordered 2D
structures of metal nanoparticles. Several approaches have become known using
particles in the size range 3–5 nm [162–168]. J. R. Heath et al. reported [163] that
success in the organization of alkanethiolate-protected nanoparticles using LB
techniques typically depends on the ratio of ligand length to surface pressure.
269
270
4 Organization of Nanoparticles
a
Fig. 4-25 Sketch of variously oriented Au55 (PPh3 )12 Cl6 clusters,
indicating the PPh3 shapes and their positionings towards each
other, resulting in square or hexagonal closed-packed structures
(a). Full space models of both orientations (b).
Ligand-stabilized Au55 clusters only ca. 2 nm in size were first organized in LB
films in 2000 [169], but only relatively small ordered areas were observed. Much
better results were obtained when the surface pressure was increased. Under these
conditions, Au55 clusters coated with PPh3 ligands, some of which had been substituted by hydrophilic derivatives to make the particles amphiphilic, formed much
larger areas of ordered particles [170]. Repeated compression of the films to 35
mN/m with alternating relaxation gave quite adequately ordered films of dimensions 0.25 mm 0.35 mm.
From the moment of introduction of the photonic bandgap concept by Jablonovitch and John [171, 172], materials possessing a three-dimensional periodicity
of the dielectric constant have attracted increasing attention. The periodically varying index of refraction in photonic crystals causes a redistribution of the density of
photonic states due to Bragg diffraction, which is associated with stop bands for
light propagation. Chemical self-assembly (bottom-up) methods have been utilized
for making 3D photonic crystals (so called artificial opals) from colloidal silica or
4.2 Metal Nanoparticles
Fig. 4-26 TEM image of a magnified cutout of perfectly squarepacked Au55 (Ph2 PC6 H4 SO3 H)12 Cl6 clusters.
Fig. 4-27 SEM image of an inverse colloidal crystal made from
PtFe nanocrystals. Courtesy of Dr. N. Gaponik.
micron-size polymer spheres. Using artificial opals as templates, so-called ‘‘inverse
opals’’ possessing pronounced photonic bandgap can be prepared by impregnation of voids with different materials followed by removal of the spheres by annealing or dissolution, which can give a complete photonic bandgap [173]. Magnetic nanocrystals can be used for the fabrication of metal-dielectric inverse opals
addressable through an external magnetic field. Figure 4-27 shows an SEM image
of an inverse FePt colloidal crystal obtained by complete impregnation of a 3D
template of monodisperse latex microspheres with FePt nanocrystal followed by
dissolution of polymer spheres in toluene.
Self-assembled structures are quite suitable for the study of the electronic behavior of ordered monomolecular films and have contributed considerably to a
271
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4 Organization of Nanoparticles
better understanding of the quantum size nature of metal nanoparticles. On the
other hand, it is quite clear that metal particles in future nanoelectronic devices
should be in structures in which the quantum dots are not self-arranged, but are
placed at definite positions. This is a big problem that will only be solved by techniques other than self-assembling.
The use of appropriate templates is one possible way to succeed; however, the
generation of the templates needs to be developed. Remarkable progress in the
organization of metal nanoparticles to form given structures involved the use of
micelles. M. Möller et al. have contributed to this development in numerous ways.
They use micellar structures in block copolymers as nano reaction vessels to generate and to organize metal nanoparticles. Although this technique is not based on
the self-assembling of metal nanoparticles, it is linked with the organization of the
micelles in the polymers. This technique differs from the self-assembly processes
described above, but it cannot be considered to solve the problem of creating artificial architectures. However, compared with the classical self-assembly procedures,
the application of preformed micelles opens up novel routes for handling nanoparticles.
A–B diblock copolymers can be used to form micelles which allow the preparation of thin, coherent films of well-developed structure [174–180]. Poly(styrene)block (b)-poly(ethylene oxide) as well as poly(styrene)-b-poly(2-vinylpyridine)
turned out to act as useful combinations for generating well-ordered compartments in which metal nanoparticles can be fabricated in various ways. Figure 4-28a
illustrates the formation of diblock copolymers from polymer A and polymer B
ending with nanospheres which can be transferred to films of highly ordered
micelles. Figure 4-28b shows schematically the formation of metal nanoparticles
inside the micelles.
As a typical example, the formation of gold nanoparticles will be described [181].
Both of the above-mentioned block copolymers can be treated with HAuCl 4 or
LiAuCl 4 . The tetrachloroaurates penetrate into the micelles in nonpolar solution to
form thermodynamically stable dispersions inside the holes. The solubilization can
be supported by ultrasound irradiation. Subsequent reduction of the gold salts results in formation of nanoparticles. Suitable reducers are hydrazine and NaBH4 .
The dispersion of the generated nanoparticles can be controlled by the domain
structure of the block copolymer [175, 176]. Because of the reduction mechanisms it is not possible to prepare one single gold monocrystal per micelle by this
method, but instead multiple nanoparticle formation in each domain is observed.
Figure 4-29 shows a TEM image of an ordered array of micelles, each loaded with
several gold particles [176]. One single nanoparticle per micelle can be obtained,
even with systematic variation of particle size, if the process is performed at the
glass transition temperature (90 C) of the poly(ethylene oxide). Particles of 6, 4,
and 2.5 nm diameter have been prepared by variable loading ratios and micelle
size and are shown in Figures 4-30a, b, and c.
Self-assembling processes leading to film formation usually occur by evaporation
of the solvent from solutions containing loaded micelles. Micellar film formation
can also be brought about by dipping platelets of different materials such as GaAs,
4.2 Metal Nanoparticles
a
b
Formation of diblock copolymers from polymer A
and polymer B, resulting in micelles (a). Formation of metal
nanoparticles inside the micelles (b).
Fig. 4-28
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4 Organization of Nanoparticles
Fig. 4-29
Gold particles inside of an ordered array of micelles.
Fig. 4-30 TEM images of gold nanoparticles of various sizes
generated in micelles of different sizes and by different loading
ratios, a 6 nm, b 4 nm, and c 2.5 nm.
InP, glass or carbon-coated copper grids into a corresponding solution and withdrawing them [182].
So, using the four kinetically controlled steps, namely reduction, mineralization,
coagulation, and film formation, the structure of nanoparticle-loaded block copolymer composites becomes available, and further use of these materials can be expected.
Gold is by no means the only material that can be incorporated into block copolymer micelles. Other metals such as palladium, silver, and cobalt have been
successfully applied, and even nanoparticles of CdS and PbS have been generated
inside micelles [183].
An interesting and important continuation of this development is the possibility
of removing the diblock copolymer material without losing the ordering of the
nanoparticles, so generating nanostructures on different surfaces and with varying
4.2 Metal Nanoparticles
inter-particle distances. Möller et al. developed an elegant method to routinely
generate these materials by the use of an oxygen plasma [184–186].
If ordered micellar films on glass or mica, containing either gold nanoparticles
or even the non-reduced tetrachloroaurates, are treated by an oxygen plasma at
200 W for 20 min, bare nanoparticles are deposited on the surfaces without any
loss of the original order. The formation of nanoparticles from AuCl 4 during the
plasma process is traced back to the existence of intermediate oxidation products of
the polymer, like CO, and also free electrons. XPS experiments prove the absence
of any residual polymer. The as-prepared particles are strongly fixed on the surfaces
and cannot be removed by rinsing with solvents or by rubbing with a soft tissue.
Another remarkable advantage is associated with this method of depositing metal
nanoparticles on surfaces: the particle size, the inter-particle distances and thereby
the particle density can be varied, and additionally the extension of the covered
surfaces is in principle not limited. The three examples in Figure 4-31 are cutouts
of 3 cm 3 cm areas and demonstrate the versatility of this method.
The important point is to use diblock copolymers of different molecular weights
and different gold contents. A poly[styrene(800)-b-2-vinylpyridine(HAuCl 4 )0:5 (860)]
(number of monomer units in brackets) gives gold colloids 12 nm in height (AFM)
and a periodicity of 80 nm (Figure 4-31a). A styrene:vinylpyridine ratio of 325:75
with the same equivalent of HAuCl 4 results in 2 nm particles with 30 nm distances (Figure 4-31b), whereas a change to 1700:450 with only 0.1 equivalent of
gold acid gives 1 nm particles of 140 nm periodicity (Figure 4-31c) [184].
Other diblock copolymers and metal compounds such as H2 PtCl6 , Pd(Ac)2 ,
TiCl 4 or FeCl3 have been successfully used. The less noble metals Ti and Fe are
deposited as oxides after plasma treatment.
Gold nanoparticles containing films of diblock copolymers can be used for
nanolithographic structuring of surfaces and so may contribute to the search for
techniques that allow routine fabrication of nanostructured surfaces. Semiconductor substrates such as GaAs, InP, or layered systems can be covered by micellar
films containing gold nanoparticles. If such composites are etched by an Arþ
beam, nanostructured surfaces are generated owing to the fact that the nanoparticulated gold is much more rapidly etched than bulk gold or than pure poly(styrene) and poly(vinylpyridine), so that all positions containing gold can be
transferred into holes in the underlying semiconductor, as illustrated in Figure
4-32.
In contrast, if a hydrogen plasma is used instead of argon, all organic material is
removed first, so that only the metal nanoparticles remain. On subsequent etching
by an ion beam chlorine plasma, GaAs cylinders of up to 80 nm in height could be
prepared [187]. Gold nanoparticles, generated from a gold-loaded diblock copolymer film by an oxygen plasma, also act as a mask during continued etching by
oxygen [188]. As a result, gold-capped diamond pillars with a diameter roughly
corresponding to the size of the gold particles are generated.
The use of so-called S-layers [189] is a combination of self-organization and
spatial patterning. S-layers consist of two-dimensional protein crystals. They are
formed by nature as the outermost cell surface layer (S-layer) of prokaryotic or-
275
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4 Organization of Nanoparticles
Fig. 4-31 TEM images of cutouts of 3 3 cm 2 areas of gold
nanoparticles produced in diblock copolymers of different
molecular weights and different gold contents.
ganisms. The subunits can recrystallize into nanoporous monolayers in suspension, at liquid-surface interfaces, on lipid films, or on solid substrates. S-layers of
Bacillus sphaericus CCM 2177 have been used to generate ordered arrays of 4–5 nm
gold particles, with a 13.1 nm repeat distance, from AuCl 4 ions [190]. Spontaneous self-assembly processes of 5 nm gold nanoparticles occur at the S-layer of
4.2 Metal Nanoparticles
Fig. 4-32 Scheme of etching semiconductor substrates starting
with gold nanoparticles containing micellar films and followed
by Arþ beam etching to remove the gold nanoparticles. This
forms holes in the substrate with the same diameter as that of
the gold nanoparticles.
Deinococcus radiodurans to give micron-sized ordered domains [191]. Arrays of 1.9
nm platinum particles were achieved from Pt salts in the S-layer of Sporosarcina
ureae [191]. These are of square symmetry with a lattice constant of 13.2 nm.
Aimed Structures
Up to now we have discussed only two-dimensional arrangements of metal nanoparticles that were based on self-assembly mechanisms, although, as in the case of
micellar films, some influence on particle size and periodicity was possible. Arrangements of nanoparticles that are completely free of natural self-assembly processes, thus allowing the formation of arbitrary arrays, seems possible, since Mirkin et al. developed so-called dip-pen nanolithography [192, 193]. This technique is
based on the transport of molecules from an atomic force microscope tip through a
water meniscus onto a flat surface that attracts the molecules by chemical bonding.
The water meniscus is naturally formed in the laboratory’s atmosphere. So, moving the tip over the surface, the corresponding molecules are transported to the
surface to form programmed patterns. Gold surfaces are very much suited to be
decorated with thiol molecules from the tip. The general method of depositing
regularly oriented nanoparticles on a surface is shown in Figure 4-33 [194].
4.2.2.3
277
278
4 Organization of Nanoparticles
Fig. 4-33 Formal illustration of dip-pen
nanolithography. An AFM tip is loaded by thiol
molecules, here 16-thiohexadecanoic acid.
Because of the presence of a ‘‘natural’’ water
meniscus, the thiols are transported onto a
gold surface via the moving tip (a) to form
any kind of pattern (b). Passivation of the
unpatterned regions by alkylthiols (c) enables
the deposition of nanopheres with basically
modified surfaces by interactions with the
carboxylic groups (d).
An array of dots of molecules that will attract the nanoparticles in a second step
are deposited on the substrate. 16-thiohexadecanoic acid has been used to bind
positively charged particles such as protonated amine or amidine-modified polystyrene spheres as building blocks from solution. Beforehand, the unpatterned regions of the gold substrate were passivated by an alkanethiol.
Orthogonal gold assemblies have been fabricated by varying this method in such
a way that, after patterning of the gold surface with 16-thiohexadecanoic acid and
passivation of the unpatterned regions with 1-octadecanethiol, alkylamine-modified
oligonucleotide (a ¼ TCTCAACTCGTAA10 ) is then chemically bound to the carboxylic groups. Several steps of rinsing with buffer solutions, washing with deionized water, and drying follow. The superiority of the dip-pen technique is
impressively demonstrated by adding a second orthogonal structure to the first
one: it is possible to replace 1-octadecanethiol by 16-thiohexadecanoic acid molecules coming from the AFM tip under usual writing conditions (0.5 nN, 55%
humidity, 23 C). A second, but different, alkylamine-modified oligonucleotide
(b ¼ A10 CGCATTCAGGAT) is then combined with the acid functions, as in the
first step. So, two nanopatterns, consisting of two different oligonucleotide sequences, can be created. The a,b-modified substrate is then treated with a 0 b 0 ¼
TACGAGTTGAGAATCCTGAATGCG with a 0 and b 0 complementary to a and b.
The deposition of gold nanoparticles of two different sizes happens by adding aand b-substituted 13 nm and 30 nm particles. The complement of a and a 0 and b
4.2 Metal Nanoparticles
a
b
a Schematic illustration of the formation of goldnanostructures via complementary oligonucleotides. b AFM
image of as-prepared patterns consisting of gold nanoparticles
of different sizes.
Fig. 4-34
and b 0 respectively ends up with a regular pattern of both types of nanoparticles
[195]. Of course, the fabrication of 2D arrangements of only one type of particle is
possible as well. Figure 4-34a illustrates schematically the formation of the nanostructures via complementary oligonucleotides. Figure 4-35b shows an AFM topography image of different Au particles after the orthogonal assembly process.
It can be foreseen that not only can two different sizes of the same kind of metal
be used to generate these or other structures, but that different metals can be used
as well, opening up manifold possibilities with respect to future applications.
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4 Organization of Nanoparticles
At this point we can conclude that the two-dimensional organization of metal
nanoparticles has developed very promisingly. From very simple self-organization
via guided self-assembling to the controlled patterning of surfaces, there is a wide
variety of routes available now that should in principle make applications in nanoelectronics possible. It should be stated, however, that not all of the organized particles mentioned above can really be used as quantum dots since they are still too
big to exhibit those very special electronic properties that are needed for singleelectron transitions at room temperature. The quantum-sized characteristics are
discussed in detail in Section 5.2.
4.2.3
One-dimensional Arrangements
The expression ‘‘one-dimensional’’ is not very well defined. In principle, onedimensionality could be reduced to the atomic scale, simply considering atoms as
the smallest building blocks of wires. In practice this very stringent definition
would not make much sense. Nanoparticles as building blocks arranged in only
one direction justifies the term ‘‘one-dimensional’’ just as well. However, how do
we consider structures of two, three or four parallel rows of ‘‘one-dimensional’’
chains of nanoparticles? It seems reasonable to consider the aspect ratio, i.e. the
ratio of length to width of an assembly of building blocks. Three times three particles could be considered as a small two-dimensional area, whereas a structure
of three times a hundred or a thousand nanoparticles is better described as onedimensional. In some cases, if the building blocks are rather big (or are very small
three-dimensional particles) or the aspect ratio is not very well expressed, one also
speaks of quasi one-dimensional structures. Therefore, in the following, there will
no longer be a distinction between a truly one-dimensional assembly of nanoparticles and a merely quasi one-dimensional assembly.
The one-dimensional organization of metal nanoparticles can be achieved via
appropriate templates or by special techniques to be discussed below. There is
a very important difference between these and 2D arrangements, where selfassembly processes can be initiated without specific templates.
Templates to organize metal nanoparticles in one direction are rare. Steps on
crystal surfaces have occasionally been observed to adsorb nanoparticles in preference to flat terraces, so giving rise to 1D orientations. However, this route has not
yet been developed into a common and routinely applicable procedure.
Nanopores are in principle promising candidates if they can be filled with
nanoparticles in such a way that the resulting composites can be used for planned
investigations or applications. In this connection, nanoporous alumina plays a
certain role, since it can easily be prepared as thin films equipped with 10 9 –10 11
pores per cm 2 , all running parallel through the membrane. These films are
formed by anodic oxidation of aluminum surfaces in polyprotic aqueous acids, the
applied voltage determining the pore width. The pore length simply depends on
the anodization time. So, pores from about 10 up to 200 nm in width and up to
some hundreds of mm in length are available [196–200]. Indeed, several attempts
4.2 Metal Nanoparticles
Fig. 4-35 A 7 nm pore in a nanoporous alumina membrane
filled with Au55 gold clusters, imaged by TEM after sectioning.
have been made to fill these pores with metal nanoparticles [196, 201–203]. Preformed particles can be transferred into the pores from solution by vacuum induction, or they can be generated inside the pores by decomposition of previously
deposited appropriate complexes. Figure 4-35 shows a TEM image of a sectioned
piece of a membrane filled with gold nanoparticles [196].
However, as it turned out, these templates did not yet lead to fully convincing
results, i.e., to the perfect wires of nanoparticles which are necessary to study
electronic conduction mechanisms, since the lack of only one particle in a row
would interrupt the electric conductance. Among other causes of this difficulty, the
main one was the length of the pores. For instance, in a pore of a 1 mm thick
membrane (a minimum thickness that could still be handled), 500 particles of size
2 nm have to be arranged without defects. This is practically impossible.
Polymer molecules could be promising alternative templates for assembling
metal nanoparticles on their surfaces. DNA is one of several candidates for assemblies of up to several microns [204]. Ligand-protected platinum particles of 2–3
nm do in fact add to natural DNA in buffer solution in a very regular fashion and
can easily be observed by TEM [211]. Figure 4-36 shows an assembly of DNA
strands covered with Pt nanoparticles. These results indicate that DNA might
indeed act as a suitable template.
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4 Organization of Nanoparticles
Double-stranded l-DNA decorated by 2–3 nm Pt
particles bearing a ligand shell of ()-cinchonidine and acetic
acid.
Fig. 4-36
l-DNA has also been used for covering with 3–5 nm Pd particles [206]. Here,
preformed particles were not used, but were generated on the DNA surface by
reduction of Pd(CH3 COO)2 with a mixture of sodium citrate, lactic acid, and dimethylamine borane.
Colloidal platinum particles have also been generated on DNA, starting with
an impregnation step with dichloro(2,2 0 :6 0 ,2 00 -terpyridine)platinum(II) or cisPt(NH3 )2 Cl2 , followed by reduction with NaBH4 [207]. Water-soluble thiolstabilized ca. 2 nm Au particles and l-DNA combine to form real 1D assemblies or
bundles of these filaments [208].
All these methods give quasi one-dimensional, but more or less disordered,
structures. A rather interesting route to ordered arrangements of gold nanoparticles in combination with DNA has recently become known [209]. The method is
based on Coulomb interactions between the negatively charged phosphate backbone of synthetic and natural DNA and positively charged 3.5 nm gold nanoparticles prepared from HAuCl 4 and NaBH4 and capped with lysine molecules.
TEM investigations of these hybrid systems indeed prove the formation of equidistant rows of nanoparticles between parallel DNA sequences, as sketched in
Figure 4-37.
4.2 Metal Nanoparticles
Fig. 4-37 Sketch of an arrangement of DNA/Au nanoparticles.
Positively charged 3.5 nm gold nanoparticles, capped with
lysine, add to the negatively charged phosphate backbones of
DNA to build up equidistant rows.
A unique kind of one-dimensionally organized metal nanoparticles in connection with DNA as templating material was observed when l-DNA was reacted with
Au55 (Ph2 PC6 H4 SO3 H)12 Cl6 [210]. As it turned out, these clusters do not only interact with DNA, like the above-mentioned 2–3 nm Pt particles, by assembling
along the double-stranded DNA, but occupy the major grooves of B-DNA, the preferred DNA version in aqueous solution. Molecular dynamics simulations indicated that bare Au55 clusters 1.4 nm in diameter fit perfectly into the grooves.
The loss of parts or even all of the original phosphine ligands can be understood as
a substitution reaction of phosphines by phosphates of the DNA backbone, acting
as a kind of polydentate system. This can be seen from Figure 4-38a. However, the
surprising effect is the degradation of these 1.4 nm particles in the B-DNA major
grooves to 0.6–0.7 nm particles in ultra-high vacuum, probably due to the shrinkage of the major grooves of B-DNA, which changes to A-DNA under these conditions. Again, the grooves of this type of DNA fit perfectly to the observed size
of particles, which correspond reasonably with Au13 . This situation is shown in
Figure 4-38b. Moreover, these Au13 -DNA hybrid systems organize in such a way
that, probably via gold-gold interactions, equidistant wires of nanoparticles result.
The ordering parameter is without any doubt the A-DNA. The simulation and the
experimental result are shown in Figure 4-38c.
The formation of the nanowires obviously occurs during drying on the grid used
for transmission electron microscopy. Au55 clusters have never been observed to
degrade under comparable conditions. This is obviously a unique instance of a
change of a molecular conformation leading to a controlled degradation of nanoparticles. Under the conditions of electron beam irradiation, however, the clusterDNA hybrid system is rather unstable. In the course of seconds, the lines break
283
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4 Organization of Nanoparticles
Fig. 4-38 a Molecular modeling resulting
from the combination of B-DNA with bare
Au55 clusters trapped in the major groves.
b Transition from B- to A-DNA, linked with
shrinking of the major groves from 1.4 to
0.7 nm causes the degradation of the Au55
clusters to 0.6–0.7 nm particles. These hybrid
systems then combine to give equidistant
wires (c).
4.2 Metal Nanoparticles
Fig. 4-38
(continued)
down to form larger gold particles, the formation of which can be observed under
the microscope. It is not clear whether a beam-conditioned decomposition of DNA,
of the wires or of both causes the degradation. Figure 4-39b shows the situation
after the decomposition of the same area shown in Figure 4-39a.
One-dimensional arrangements of metal nanoparticles using templates other
than DNA have also been generated successfully. Polymer molecules were used in
two cases for decoration with Au55 (PPh3 )12 Cl6 clusters. Substituted poly( p-phenyleneethynylenes) turned out to readily chemisorb the clusters, probably owing to pp interactions between the electron-rich polymer and the phenyl groups in the
clusters’ ligand shells. In the special case of the I-octyl-substituted derivative (PPEI-oct), small bundles of polymer molecules can be generated in very dilute solutions [211]. PPE-I-oct (0.5 mg per liter dichloromethane) forms bundles of diameter
ca. 10–12 nm and length 250 nm. These thin chains of polymers are able to adsorb
the clusters in a truly one-dimensional manner, as can be seen from Figure 4-40.
Another polymer suitable for collecting the same nanoparticles is poly(vinylpyrrolidone) (PVP). Distributed in a thin dichloromethane film on a water surface
in a Langmuir-Blodgett trough, cluster-covered individual polymer molecules can
be imaged by AFM. However, when surface pressure is increased, the PVPnanoparticle wires continuously form dense filaments, which show numerous
knots, leading to a two-dimensional network [212]. It is assumed that these knots
consist of gold nanoparticles, since PVP, imaged under the same conditions, does
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286
4 Organization of Nanoparticles
Fig. 4-39 a TEM image of gold nanowires consisting of 0.6–
0.7 nm particles (see magnified cutout) arranged equidistantly
(0.5 nm). b The same area after continued e-beam irradiation,
showing the formation of large gold nanoparticles from the
wires.
4.2 Metal Nanoparticles
Fig. 4-40 TEM image of one-dimensional arrangements
of Au55 (PPh3 )12 Cl6 clusters along chains of
poly( p-phenyleneethynylenes).
not exhibit corresponding behavior. Figure 4-41 shows the formation of these filaments, which can be transferred to solid substrates without degradation because of
the linkage via nanoparticles.
Arrangements like these could be of interest with respect to future nanoelectronic devices using Au55 clusters as quantum dots.
A template-free fabrication of one-dimensionally organized Au55 clusters using a
controlled degradation of ordered monolayers has been found. As described in
Section 4.2.2.1, islands of square and hexagonal structures of Au55 (PPh3 )12 Cl6
clusters are formed on the surface of water, which was previously covered with a
thin film of clusters in dichloromethane, when the solvent is evaporated. If these
islands are transferred to solid substrates under well-defined conditions, they are
degraded into strictly parallel stripes, each consisting of 3–4 rows of clusters [213].
The transfer conditions found to work best consist of a 20 withdrawal angle and a
speed of 10 cm min1 . The formation of these patterns is explained by the oscillation of the water meniscus at the water-substrate boundary. The complex mechanism is illustrated in Figure 4-42. The monolayer is divided parallel to the transfer
direction. Figure 4-43 shows a TEM image of an as-prepared pattern of stripes of
Au55 clusters.
This kind of organization of nanoparticles is related to a method based on wetting instabilities of monomolecular layers when transferred onto solid substrates.
Low-pressure monolayers of l-a-dipalmitoyl-phosphatidylcholine (DPPC) on water
degrade, forming a pattern with parallel hydrophilic channels 200–300 nm in
287
288
4 Organization of Nanoparticles
Fig. 4-41 a AFM image of Au55 (PPh3 )12 Cl6 clusters decorating
chains of poly(vinylpyrrolidone) forming a 2D network.
b A magnified cutout.
4.2 Metal Nanoparticles
Fig. 4-42 Sketch of the mechanism generating stripes of
Au55 (PPh3 )12 Cl6 chains by the division of ordered islands into
sections caused by the oscillating meniscus at the watersubstrate boundary.
width when transferred to mica at a speed of 1 mm s1 [214]. The channels of
such structures can be used to deposit nanoparticles in one direction. Thiolstabilized Au55 clusters, dissolved in 1-phenyloctane, were used to fill the channels,
aggregates being formed during removal of the solvent. The DPPC areas are almost completely free of clusters owing to their nonwettability.
Because of the magnetic dipolar interactions, one-dimensional self-assembly is
rather typical for magnetic nanoparticles. For a magnetic single-domain nanoparticle of radius r, the magnetic dipole moment arising from the alignment of
electron spins is m ¼ 4pr 3 Ms =3, where Ms is the saturation magnetization of the
respective bulk material. The dipole-dipole interaction between two magnetic particles at contact is proportional to m 2 =s 3 z r 6 =s 3 , here s is the effective hard
sphere diameter, consisting of the magnetic core diameter and the thickness of
289
290
4 Organization of Nanoparticles
Fig. 4-43 TEM image of equidistant stripes of Au55 (PPh3 )12 Cl6
clusters consisting of 3–4 individual 1D chains.
the surfactant layer [215]. These dipolar interactions favoring the head-to-tail arrangement of dipoles can compete with nondirectional van der Waals forces and
steric repulsion, inducing anisotropic agglomeration of nanoparticles. In 1970, de
Gennes and Pincus [216] theoretically predicted that colloidal particles with a
magnetic dipole moment should be able to minimize the magnetostatic energy via
self-assembly into flexible chains, as schematically shown in Figure 4-44a. Simulations show that such chains can be present in colloids of superparamagnetic
particles even in zero magnetic field [217, 218]. Only in 2003 did Butter et al. [219]
provide direct in situ observation of such linear chains of particles present in a
dispersion of iron nanoparticles at zero field. They used a cryogenic TEM technique based on fast vitrification of the colloidal solution followed by low-dose TEM
imaging of metal nanoparticles inside the vitrified film. A clear transition from a
single particle population observed for 4 nm iron nanoparticles to a system of
worm-like chains for particles larger than ca. 12 nm was observed, and this was
explained by rapid increase of dipolar interaction with particle size.
Unpaired dipoles present at the ends of a linear chain of nanoparticles can pair,
yielding the closed ring shown in Figure 4-44b. Such structures were observed by
4.2 Metal Nanoparticles
a
b
c
H
Fig. 4-44 Dipolar chains formed by magnetic
particles. a A semi-flexible chain of magnetic
colloidal particles showing the head-to-tail,
paired dipoles in the middle section and the
unpaired dipoles at the ends. b A self-
assembled closed ring of magnetic particles
allows all magnetic dipoles to be paired.
c Straight chains of magnetic particles can
form and align in a magnetic field.
Tripp et al. [220] for 27 nm cobalt nanoparticles (weakly ferromagnetic at room
temperature), which self-assemble into bracelet-like rings. Typically, the rings consisted of 5–12 particles and were 50–100 nm in diameter. The key point for selective formation of the nanoparticle rings instead of large aggregates was increased
viscosity of solvent during an earlier stage of drying the colloidal solution on a
TEM grid.
An external magnetic field orients the magnetic dipoles of individual nanoparticles in one direction, facilitating the formation of linear head-to-tail superstructures. The chain of dipole-dipole interacting nanoparticles will align itself in the
direction of an externally applied magnetic field, as shown in Figure 4-44c. Another striking effect of the magnetic field is the dependence of the lateral width of
1D-nanoparticle aggregates on the strength of the applied field, as was demonstrated by Giersig et al. using the example of cobalt nanoparticles (Figure 4-45)
[132]. Progressive increase in the width of nanoparticle chains with increasing
magnetic field can be explained by the attractive chain-chain interactions between
parallel dipolar chains of particles aligned by the magnetic field [219, 221]. In the
case of a substrate positioned vertically, the self-assembly of magnetic nanoparticles into 1D chains of variable width has been explained as a cooperative effect of
the external magnetic field and the force of gravity [222].
The shape of the magnetic nanoparticles can also affect the preferential dimensionality of self-assembled superstructures. Thus, hcp-Co nanodisks synthesized by
Puntes et al. show a clear predisposition to stack face-to-face and remain perpendicular to the substrate, forming long ribbons (Figure 4-45) [223]. In the TEM
images, these resemble ribbons of nanorods. However, tilting of the TEM grid
shows that the standing disks vary in thickness when tilted in the short direction
291
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4 Organization of Nanoparticles
0.1
0.35
0.3
0.8
Tesla
Fig. 4-45 TEM images of 1D ordered cobalt nanoparticles
self-assembled in magnetic fields of different strengths.
Reproduced from [132], Copyright 2002, with permission
from Elsevier Science.
Fig. 4-46 TEM picture of stacked face-to-face
hcp-Co nanodisks. The pictures on the left
correspond to no tilt and those on the right to
a 25 tilt. The bar is 200 nm. Tilting direction
corresponds with the double arrows.
Reproduced from [223], Copyright 2002, with
permission from the American Chemical
Society.
and remain of constant length when tilted in the long direction (Figure 4-46), verifying the disk shape.
More or less all attempts to organize nanoparticles are driven by their possible
use in future nanoelectronic devices, above all as miniaturized transistors. This
highly desirable goal can only be reached if we succeed in generating structures
4.2 Metal Nanoparticles
Scheme 4-1 Process of formation of light-induced positions to
add 2.6 nm gold nanoparticles by ligand exchange (a–c).
Additional gold nanoparticles can be coordinated via the
removal by dithiols of more weakly bound amines.
which are not made by nature, but instead follow complex designs. The precise
placement of nanoparticles, acting as single electron transistors, is therefore the
ultimate dream of any research in that field.
The very first attempts have been made. The present state of the art is based on
the use of appropriately nanostructured surfaces that are chemically modified in
such a way as to irreversibly bind nanoparticles onto the structure. One possible
route has been pointed out by J. R. Heath et al. However, it still leads to rather
broad structures [224]. They used light-controlled arrangement of gold nanoparticles on glass and silicon substrates. These are first covered with a monolayer of
(3-aminopropyl)ethoxydimethylsilane, and this is followed by the addition of nitroveratryloxycarbonyl glycine (NVOC-GLY) to yield a surface with light-sensitive
NVOC functions. Exposure to UV/Vis light through a mask results in a pattern of
free and protected amino groups. Gold particles (2.6 nm), protected by a shell of
1-dodecylamine, are then used for addition to the free amino groups. Ligand exchange reactions between the 1-dodecylamine and the free amine groups of the
surface lead to the decoration of the exposed areas with gold nanoparticles. The
process is illustrated in Scheme 4-1. As can also be seen from this figure, additional particles can be placed in successive steps via dithiol ligands removing more
weakly bound amines.
293
294
4 Organization of Nanoparticles
Scheme 4-2 Generation of Au55 patterns on a
prestructured surface. An alkyl-terminated selfassembled monolayer (a) is oxidized along
traces by a metalized AFM tip inducing electric
pulses (b). Alkene-terminated silanes are then
added specifically to the COOH functions (c).
This is followed by thiolization of the CbC
double bonds (d) and addition of water-soluble
Au55 (Ph2 PC6 H4 SO3 Na)12 Cl6 clusters (e).
4.2 Metal Nanoparticles
The best existing tool for fabricating nanostructured surfaces is, without doubt,
the AFM tip. Mirkin et al. used it to develop so-called dip-pen nanolithography to
deposit thiol molecules on gold surfaces [192, 193].
Conductive AFM tips can be used to emit electrical pulses to affect the top
functions of ordered monolayers [225–227]. Here, the tip operates as a nanoelectrical pen, inscribing chemical information into appropriate surfaces. These traces,
chemically modified, can then be used to deposit nanoparticles for further applications.
As can be seen from Scheme 4-2, a C18 alkyl monolayer on silicon has been oxidized by a tip-induced nanoelectrical reaction, generating carboxylic functions
[228]. A second self-assembled monolayer, produced by nonadecenyltrichlorosilane, on top of the COOH groups functionalizes the carboxylic tracks by CbC
Fig. 4-47 a Lines of SH functionalized traces corresponding to
step d in Scheme 4-2 and height profile. b Au55 -decorated lines
corresponding to step e in Scheme 4-2. The height profile
agrees well with the addition of ca. 2.5 nm to the 2 nm height
of the thiol-functionalized traces.
295
296
4 Organization of Nanoparticles
double bonds that can be attacked photolytically by H2 S to give the corresponding
thiol-functionalized pattern. The purpose of using BH3 is to reduce SaS bonds,
possibly formed during the process. The water-soluble gold cluster Au55 (Ph2 PC6 H4 SO3 Na)12 Cl6 was used for deposition on the thiolated traces. The structures are
remarkably thermally stable, as was shown by annealing up to 120 C.
With this technique, J. Sagiv et al. generated a series of planned patterns of 1.4
nm gold particles, whose electronic properties are known to make them promising
candidates for future single-electron switches. Figures 4-47 to 4-50 show some of
these structures. AFM analyses of the patterns indicate the existence of singlecluster layers. Figure 4-47 demonstrates a pattern of parallel wires and the corresponding distance-height profiles. The wires in (a) consist of the bilayer system
only (step d in Scheme 4-2), whereas (b) images the same pattern after the addition
of the gold clusters (step e in Scheme 4-2). As can be seen from the distance-height
profiles, there is an increase in height from ca. 2 nm in (a) to ca. 5 nm in (b), corresponding to a particle height of ca. 3 nm. Considering the inherent limitations of
AFM height determination, this value agrees quite well with the cluster size of ca.
2.5 nm.
Details cannot be distinguished in Figure 4-47. However, in the magnified cutouts in Figure 4-48, gold clusters placed in two densely packed parallel rows can be
observed. The clusters seem to be arranged along the periphery of the tracks.
Figure 4-49 shows a pattern of dots, each consisting of assemblies of 2–3 partially resolved clusters.
Fig. 4-48 Magnified pieces of Au55 cluster-decorated lines of
Figure 4-47, indicating 2 rows of clusters per line.
4.2 Metal Nanoparticles
Fig. 4-49 AFM images of a regular arrangement of dots of
Au55 clusters, each dot consisting of 2–3 clusters.
The capabilities of this method, provided that the experimental conditions are
optimized, can be seen from Figure 4-50. Ideal one-dimensional wires of individual
particles become available, and it would even be possible to place single clusters in
predefined intrawire gaps.
So far, this route represents the most promising recipe for the creation of
planned architectures of quantum dots. Of course, there is still a long way to go
before practical applications become realizable. On the other hand, there exists no
297
298
4 Organization of Nanoparticles
Fig. 4-50 Optimized pattern of ideal 1D cluster wires and
single clusters in predefined intrawire gaps.
other system which could be built up from particles as small as the 1.4 nm gold
particles mentioned above and which is electronically exposed in such a manner
that single electrons can be used to switch, even at room temperature. This vision
justifies any efforts to further improve the techniques to build up working constructions.
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305
5
Properties
5.1
Semiconductor Nanoparticles
5.1.1
Optical and Electronic Properties of III–V and II–VI Nanoparticles
Uri Banin and Oded Millo
Introduction
Semiconductor nanocrystals are novel materials with properties lying between the
molecular and solid-state regimes, these properties, uniquely, being controlled by
size and shape [1–8]. Containing hundreds to thousands of atoms, 20–200 Å in
diameter, nanocrystals maintain a crystalline core with the periodicity of the bulk
semiconductor. However, as the wave functions of electrons and holes are confined
by the physical nanometric dimensions of the nanocrystals, the structure of the
electronic levels and the resultant optical and electrical properties are greatly
modified. Upon reducing the semiconductor size to the nanocrystal regime, a
characteristic blue shift of the band gap appears, and a discrete level structure
develops as a result of the ‘‘quantum size effect’’ in these quantum structures [9].
In addition, because of their small size, the charging energy associated with the
addition or removal of a single electron is very high, leading to pronounced singleelectron tunneling effects [10–12]. These phenomena provide ample scope for
basic scientific research, which has attracted and continues to attract considerable
attention. Because of their unique optical and electrical properties, nanocrystals are
likely to play a key role in the emerging new field of nanotechnology in applications ranging from lasers [13, 14] and other opto-electronic devices [15–17] to biological fluorescence marking [18–20]. Here we discuss the optical and electronic
properties of colloidally grown quantum dots (QDs) and quantum rods (QRs). In
the present chapter, particular attention is focused on the intuitive and appealing
‘‘artificial atom’’ picture of the electronic structure of semiconductor nanocrystals.
In this respect, the combination of tunneling and optical spectroscopy methods
has proven to be a powerful approach to the study of size- and shape-dependent
level structure and single-electron charging in these systems.
Colloidally grown nanocrystals have several favorable features such as con5.1.1.1
Nanoparticles: From Theory to Application. Edited by Günter Schmid
Copyright 8 2004 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
ISBN: 3-527-30507-6
306
5 Properties
tinuous size control and chemical accessibility resulting from their overcoating
with organic ligands. This chemical compatibility enables powerful chemical or
biochemical means to be used to assemble nanocrystals in a controlled manner
[21–27]. Further powerful control, of optical and electronic properties of nanocrystals has been achieved using colloidal synthesis. Heterostructured nanocrystals
consisting of semiconductor shells grown on a core were developed [28–36]. Enhanced fluorescence and increased stability can be achieved in these particles
compared with cores overcoated by organic ligands. Shape control was also achieved by suitable modification of the synthesis, allowing the preparation of rodshaped particles – quantum rods (QRs) [6, 37, 38]. These manifest the transition
from 0-D quantum dots to 1-D quantum wires in the sense that strong quantum
confinement holds for the radial direction but weak confinement for the axial
direction [39, 40]. QRs exhibit electronic and optical properties differing from
those of QDs, e.g., polarized emission [41], which leads to polarized lasing [14].
Additionally, the rod geometry was found to be advantageous in polymer nanocrystal-based solar cells, because of improvement in charge transport properties
[42].
In this chapter we deal with the electronic structure and the optical and electrical
properties of colloidally grown semiconductor nanocrystals, and concentrate on the
application of tunneling and optical spectroscopy to this problem. In Section
5.1.1.2, we describe the application of optical and tunneling spectroscopy to the
study of semiconductor nanocrystals, using the artificial atom analogy to represent
their level structure. In Section 5.1.1.3, the modifications associated with core/shell
nanocrystals are discussed, and, in particular, the atomic-like envelope wave functions are visualized using scanning tunneling microscopy. Section 5.1.1.4 discusses
the electronic properties of quantum rods, and Section 5.1.1.5 describes recent
results on optical gain and lasing obtained with semiconductor nanocrystals.
Semiconductor Nanocrystals as Artificial Atoms
From the early work on the quantum confinement effect in colloidal semiconductor QDs, electronic levels have been assigned according to the spherical symmetry
of the electron and hole envelope functions [9, 43]. The simplistic ‘‘artificial atom’’
model of a particle in a spherical box predicts discrete states with atomic-like state
symmetries, e.g., s and p. To probe the electronic structure of II–VI and III–V
semiconductor nanocrystals, a variety of ‘‘size-selective’’ optical techniques have
been used, mapping the size dependence of dipole allowed transitions [44–48].
Theoretical models based on an effective mass approach with varying degree of
complexity [45, 49] as well as pseudopotentials [50, 51] were used to assign the
levels. A model system that was studied extensively in this context is that of CdSe
nanocrystals, for which spectroscopic mapping of electronic states using PLE
(photo luminescence excitation) was performed by Bawendi and coworkers [44,
45]. The size-selective PLE method was also utilized to examine the size-dependent
level structure of InAs QDs, an additional system that has attracted interest [47,
48]. InAs is a narrow-gap semiconductor (Eg ¼ 0:418 eV) with a large Bohr radius
a 0 of 340 Å, whereas CdSe has a significantly larger band gap, Eg ¼ 1:84 eV and
5.1.1.2
5.1 Semiconductor Nanoparticles
photoluminescence
(A)
307
(B)
Edet =1.18 eV
a
14 Å
c
PLE signal
PL intensity [a.u.]
T=10 K
b
b
1
0.8
1.2
6 7
8
9
b
a
3 4
1.6
2.0
Energy (eV)
Fig. 5-1 A Optical spectroscopy of InAs
nanocrystals with a mean radius of 2.5 nm.
The top frame shows the absorption (solid
line) and the PL (dotted line) for the sample.
The lower frame shows a size-selected PLE
spectrum, where eight transitions are resolved,
measured with a narrow detection window
24 Å
c
a
5
20 Å
c
a
2.4
1
b
32 Å
c
1.4
1.8
2.2
Energy [eV]
positioned as indicated by the arrow in the top
panel. The weak transition E2 is not detected
in this QD size. B Size-dependent PLE spectra
for four representative InAs QD radii. The band
gap transition (a), and two strong excited
transitions (b and c) are indicated.
smaller Bohr radius (a 0 ¼ 55 Å). Therefore both systems serve as a prototypical
framework for the study of quantum confinement effects.
Figure 5-1A shows the typical spectral features measured at 10 K for a sample of
InAs nanocrystals with a mean radius of 2.5 nm. The absorption onset exhibits a
@0.8 eV blue shift from the bulk band gap. A pronounced first peak and several
features at higher energies are observed in the absorption spectra. Band edge
luminescence is observed, with no significant red-shifted (deep-trapped) emission.
In the PLE spectrum, a set of up to eight transitions are resolved. The full size dependence was measured by changing the detection window and by using different
samples, and a representative set of such PLE spectra is shown in Figure 5-1B [47].
The map of excited transitions for InAs nanocrystals was extracted from such
PLE data, and the levels were assigned using an eight-band effective mass model
[47]. Briefly, in this approach, each electron (e), and hole (h) state, is characterized
by its parity and its total angular momentum F ¼ J þ L, where J is the Bloch band
2.6
308
5 Properties
edge angular momentum (1/2 for the conduction band (CB), 3/2 for the heavy and
light hole bands, and 1/2 for the split-off band) and L is the angular momentum
associated with the envelope function. The standard notation for the electron and
hole states is nQ F , where n is the main quantum number, and Q ¼ S; P; D; . . . ;
denotes the lowest L in the envelope wave function. The calculated level separations were found to closely reproduce the observed strong transitions.
Tunneling transport through semiconductor nanocrystals can yield complementary information on their electronic properties which cannot be probed in the
optical measurements. While, in the optical spectra, allowed valence band (VB) to
conduction band (CB) transitions are detected, in tunneling spectroscopy the CB
and VB states can be separately probed. In addition, the tunneling spectra may
show effects of single-electron charging of the QD. Such interplay between singleelectron charging and resonant tunneling through the QD states can provide
unique information on the degeneracy and therefore the symmetry of the levels.
The interplay between single-electron tunneling (SET) effects and quantum size
effects in isolated nanoparticles can be experimentally observed most clearly when
the charging energy of the dot by a single electron, Ec , is comparable to the electronic level separation DEL , and both energy scales are larger than kB T [10, 52, 53].
These conditions are met by semiconductor nanocrystals in the strong quantum
confinement regime, even at room temperature, while for metallic nanoparticles,
Ec is typically much larger than DEL . SET effects are relevant to the development of
nano-scale electronic devices, such as single-electron transistors [54, 55].
However, for small colloidal nanocrystals, the task of wiring up the QD between
electrodes for transport studies is exceptionally challenging. To this end, various
mesoscopic tunnel junction configurations have been employed, such as the double barrier tunnel junction (DBTJ) geometry, where a QD is coupled via two tunnel
junctions to two macroscopic electrodes [10, 11, 56, 57]. Klein et al. achieved this
by attaching CdSe quantum dots to two lithographically prepared electrodes, and
have observed SET effects [58]. In this device, a gate voltage can be applied to
modify the transport properties. An alternative approach to achieving electrical
transport through single QDs is to use scanning probe methods. Alperson et al.
observed SET effects at room temperature in electrochemically deposited CdSe
nanocrystals using conductive atomic force microscopy (AFM) [59].
A most useful approach for tunneling-transport spectroscopy on single nanocrystals is based on using the scanning tunneling microscope (STM). For such
measurements, the nanocrystals can be linked to a gold film via hexanedithiol (DT)
molecules [12, 60], enabling the realization of a DBTJ, as shown schematically in
Figure 5-2A. An STM topographic image of an isolated InAs QD, 32 Å in radius,
is presented in Figure 5-2A (left inset). Also shown is a tunneling current-voltage
(I–V) curve that was acquired after positioning the STM tip above the QD. A region of suppressed tunneling current is observed around zero bias, followed by a
series of steps at both negative and positive bias. Figure 5-2B shows the dI=dV
versus V tunneling conductance spectrum, which is proportional to the tunneling
DOS (density of states). A series of discrete single-electron tunneling peaks is
clearly observed, where the separations are determined by both Ec (addition spec-
Tunneling Current (nA)
5.1 Semiconductor Nanoparticles
1.0
0.5
0.0
-0.5
QD
DT
-1.0
Au
A
10
dI/dV (arb. u.)
Tip
4.2 K
8
DCB+Ec
6
4
2
0
B
-2
-2
D VB+Ec
-1
Eg+Ec
Ec
0
1
2
Bias (V)
STM measurements on a single InAs
nanocrystal of radius 3.2 nm, acquired at 4.2 K.
The QDs are linked to the gold substrate by
DT molecules, as shown schematically in the
right inset. The left inset presents a 10 10
nm STM topographic image of the QD. The
tunneling I–V characteristic is shown in A, and
Fig. 5-2
the tunneling conductance spectrum in B. The
arrows depict the main energy separations: Ec
is the single-electron charging energy, Eg is the
nanocrystal band-gap, and DVB and D CB are the
spacings between levels in the valence and
conduction bands, respectively.
trum) and the discrete level spacings (excitation spectrum) of the QD. The I–V
characteristics were acquired with the tip retracted from the QD to a distance
where the bias predominantly drops on the tip-QD junction. In these conditions,
representing an asymmetric DBTJ, CB (VB) states appear at positive (negative)
sample bias, and the excitation peak separations are nearly equal to the real QD
level spacings [12, 61].
On the positive bias side of Figure 5-2B, two closely spaced peaks are observed
immediately after current onset, followed by a larger spacing and a group of six
nearly equidistant peaks. The doublet is assigned to tunneling through the lowest
CB QD state, where the spacing corresponds to the single-electron charging energy, Ec . The observed doublet is consistent with the degeneracy of the envelope
function of the first CB level, 1Se (here we revert to a simpler notation for the
CB states), which has s character. In this case, a direct relationship between the
degeneracy of a QD level and the number of addition peaks is expected. This is
309
310
5 Properties
further substantiated by the observation that the second group consists of six
peaks, corresponding to the degeneracy of the 1Pe state, spaced by Ec. This sequential level filling resembles the Aufbau principle of building up the lowest
energy electron configuration of an atom, directly demonstrating the atomic-like
nature of the QD.
The separation between the two groups of peaks is a sum of the level spacing
DCB ¼ 1Pe 1Se and the charging energy Ec. On the negative bias side, tunneling
through filled dot levels takes place, reflecting the tunneling DOS of the QD
valence band. Again, two groups of peaks are observed. The multiplicity in this
case, in contrast to the CB, cannot be clearly assigned to specific angular momentum degeneracy. In a manner similar to that described above for the CB states, we
extract a value of DVB for the level separation between the two observed VB states.
In the region of null current around zero bias, the tip and substrate Fermi energies
are located within the QD band gap where the tunneling DOS is zero, and the
energy gap Eg can be extracted from the observed spacing between the highest VB
and the lowest CB peaks. The gap region appears to be free of spectral features (see
also Figure 5-3). Upon removing the coating ligands by pyridine treatment, subgap states can be observed [62], which were attributed to surface states [63].
The tunneling conductance spectra for single InAs nanocrystals spanning a size
range of 10–35 Å in radius are presented in Figure 5-3. Two groups of peaks are
observed in the positive bias side (CB). The first is always a doublet, consistent
with the expected s symmetry of the 1Se level, while the second has a higher multiplicity of up to six, consistent with 1Pe . The separation between the two groups
and the spacing of peaks within each multiplet increase with decreasing QD
radius. This reflects quantum size effects on the nanocrystal energy levels and its
charging energy, respectively. On the negative bias side, generally two groups of
peaks, which exhibit similar quantum confinement effects, are also observed.
Here, variations in the group multiplicities between QDs of different sizes as well
as variations in the peak energy spacings within each group are observed. This
behavior is partly due to the fact that Ec is very close to DVB , the level spacing in the
valence band, as shown in Figure 5-4B. In this case, sequential single-electron
tunneling may be either addition to the same VB state or excitation with no extra
charging to the next state. An atomic analogy for this situation can be found in the
changing order of electron occupation when moving from the transition to the
noble metals within a row of the periodic table.
The comparison between tunneling and PLE data can be used to decipher the
complex QD level structure. This correlation is also important for examining possible effects of charging and tip-induced electric field in the tunneling measurements on the nanocrystal level structure. First, the size dependence of the band
gap Eg as extracted from the tunneling data can be compared with the nanocrystal
sizing curve (Figure 5-4A). The sizing curve (open diamonds) was obtained by
correlating the average nanocrystal size, measured using transmission electron
microscopy (TEM), with the excitonic band gap of the same sample [48]. To compare these data with the tunneling results, a correction term, 1:8e 2 =er, was added
to compensate for the electron-hole excitonic Coulomb interaction that is absent
5.1 Semiconductor Nanoparticles
r = 10 Å
dl/dV (arb.u.)
r = 19 Å
r = 22 Å
r = 28 Å
r = 32 Å
-3
-2
-1
0
1
2
3
Bias (V)
Size evolution of the tunneling
spectra at 4.2 K of single InAs QDs, displaced
vertically for clarity. The positions of the
centers of the zero current gap showed nonsystematic variations with respect to the zero
Fig. 5-3
bias, of the order of 0.2 eV, probably due to
variations of local offset potentials. For clarity
of presentation, the spectra are offset along
the V-direction to center them at zero bias.
Representative nanocrystal radii are shown.
in the tunneling data [9]. Agreement is good for the larger nanocrystal radii, with
increasing deviation for smaller nanocrystals. This deviation may be due to the
difference in the estimation of QD size between the experiments, which should be
more pronounced in the small-size regime [12].
Next, in Figure 5-4B, the size dependence of the higher strongly allowed optical
transitions is compared with the level spacings measured by tunneling spectroscopy. The two lower data sets (II) compare the difference between the first strong
excited optical transition and the band gap from PLE (b-a in Figure 5-1) with the
separation DVB ¼ 2VB 1VB in the tunneling data (open and full squares, respectively). The excellent correlation observed enables us to assign this first excited
transition in the PLE to a 2VB 1Se excitation, as shown schematically in the inset
of Figure 5-4A. Strong optical transitions are allowed only between electron and
hole states with the same envelope function symmetry. We thus infer that the en-
311
5 Properties
2.0
1.8
1Pe CB
1Se
Eg (eV)
1.6
1.4
I
1VB
2VB
I
II III
VB
1.2
1.0
A
0.8
35
30
25
20
15
10
Radius (Angstrom)
0.6
III
0.5
0.4
DE (eV)
312
II
0.3
0.2
0.1
0.8
B
1.0
1.2
1.4
1.6
1.8
2.0
Eg (eV)
Fig. 5-4 Correlation of optical and tunneling
spectroscopy data for InAs QDs. The inset
shows a schematic of the CB and VB level
structure and optical transitions I, II, and III
(a, b and c in Figure 5-1). A Comparison of
the size dependence of the low-temperature
optical band gap (transition I) after Coulomb
correction (open diamonds) with the band
gap measured by the STM (filled diamonds).
B Excited transitions plotted versus the band
gap for tunneling and optical spectroscopy :
The two lower data sets (II) depict the
correlation between DVB ¼ 1VB 2VB , detected
by tunneling spectroscopy (full squares), with
the difference between transition II and the
band gap transition I (open squares). The two
upper data sets (III) depict the correlation
between DCB ¼ 1Pe 1Se , determined from
the tunneling spectra (full circles), with the
difference between optical transition III and I
(open circles). Also shown is the size dependence of Ec from the tunneling data (full
triangles).
5.1 Semiconductor Nanoparticles
velope function for state 2VB should have s character, and this state can be directly
identified as the 2S3=2 valence band level.
Another important comparison is depicted by the higher pair of curves, set III.
The second strong excited optical transition relative to the band gap (c-a in Figure
5-1) is plotted along with the spacing DCB ¼ 1Pe 1Se from the tunneling spectra.
Again, excellent correlation is observed, which allows us to assign this peak in the
PLE to the 1VB 1Pe transition (Figure 5-4A, inset). The topmost VB level, 1VB ,
should thus have some p character for this transition to be allowed. From this, and
considering that the band gap optical transition 1VB 1Se is also allowed, we conclude that 1VB has mixed s and p character, in agreement with theory [47, 64]. The
good agreement between the tunneling and optical measurements indicates that
charging and the tip-induced Stark effect do not significantly perturb the QD level
structure.
Concerning the tunneling data, it is important to note that a detailed understanding of the role played by the double barrier tunnel junction (DBTJ) geometry
and the ability to control it are essential for the correct interpretation of tunneling
characteristics. In particular, the ratio between the junction capacitances determines the voltage division and therefore the observed level spacing. By varying the
tip-QD distance it is possible to modify the voltage division as discussed in [65] and
[66]. Additionally, the degree of single-electron charging effects depends on the
ratio of the tunneling rates between the two junctions. When the rate of one of the
junctions is significantly higher, charging effects are suppressed, as was achieved
by depositing InAs nanocrystals on graphite (HOPG – highly oriented pyrolytic
graphite) without linker molecules, leading to vanishing of the charging multiplets
[61], as can be seen in Figure 5-5. A similar effect was also observed for CdSe and
CdS nanocrystals by Bakkers et al. [67].
5.1.1.3 Theoretical Descriptions of the Electronic Structure
The theoretical treatments for both optical and tunneling experiments on QDs first
require the calculation of the level structure. Various approaches have been developed to treat this problem, including effective mass-based models, with various
degrees of band-mixing effects [45, 47, 49, 68] and a more atomistic approach
based on pseudo-potentials [64, 69]. Both have been successfully applied for various nanocrystal systems [50, 51, 69–72]. To model the PLE data, one has to calculate the oscillator strength of possible transitions and take into account the
electron-hole Coulomb interaction which modifies the observed (excitonic) band
gap. In the tunneling case, as discussed above, the device geometry should be
carefully modeled, and, in addition, the effects of charging and the tip-bias on the
level structure needs to be considered. The charging may affect the intrinsic level
structure and also determines the single-electron addition energy, while the tipbias may perturb the levels via the Stark effect.
Zunger and coworkers treated the effects of electron charging for a QD embedded in a homogeneous dielectric medium characterized by eout. The addition
energies and quasi-particle gap were calculated as a function of eout [73, 74]. While
this isotropic model does not represent the experimental geometry of the tunnel-
313
5 Properties
Experiment s
dl/dV (arb. u.)
(a)
p
(b)
Simulation
s
p
dl/dV (arb. u.)
314
-2
-1
0
1
2
Sample Bias (V)
Fig. 5-5 a Tunneling spectra measured on
InAs QDs of radius ca. 2 nm. The solid curve
was measured in the QD/HOPG geometry and
the dashed curve in the QD/DT/Au geometry.
b Calculated spectra showing the effect of
tunneling-rate ratio. The dashed and solid
curves were calculated with tunneling rate
ratios, G2 /G1 , ca. 1 and 10, respectively. The
image shows two QDs positioned near a step
on HOPG.
ing measurements, the authors were able to find good agreement between the
energetic positions of the peaks for several QD sizes using one value of eout . These
authors also noted that the charging energy contribution associated with the bandgap transition may be different from that within the charging multiplets in the
excited states. This difference is, however, on the order of the peak width in our
spectra.
In another approach, Niquet et al. modeled the junction parameters (capacitances and tunneling rates) and used a tight-binding model for the level structure
[75, 76]. The tunneling spectra were calculated using a rate equation method. The
authors were able to reproduce the experimental tunneling spectra, attributing part
of the tunneling peaks at negative bias to tunneling through the CB.
Atomic-like States in Core/Shell Nanocrystals – Spectroscopy and Imaging
Additional complexity and functionality are afforded for semiconductor nanocrystals through the growth of shells. Of particular interest are core/shell nanocrystals that exhibit enhanced emission quantum yield (QY) and improved stability.
In such composite core/shell structures, the shell type and shell thickness provide
further control for tailoring the optical, electronic, electrical and chemical properties of semiconductor nanocrystals. The growth of InAs/shell nanocrystals with
high QY in the technologically important near-IR spectral range is discussed in
Section 3.1.3.4 and in [34]. These nanocrystals were successfully implemented
5.1.1.4
5.1 Semiconductor Nanoparticles
s
p
1
2
dI/dV (a. u.)
InAs core
2 ML shell
6 ML shell
-2
-1
0
Sample Bias (V)
Tunneling conductance spectra of an InAs core QD
and two core/shell nanocrystals with two and six ML shells
with nominal core radii of ca. 1.7 nm. The spectra were offset
along the V direction to center the observed zero current gaps
at zero bias.
Fig. 5-6
as chromophores in polymer-nanocrystal light emitting diodes active in the telecomm range [17]. They also constitute an important system for further investigating the evolution of level structure in QDs. The combined tunneling and
optical spectroscopy approach has been further applied to investigate core/shell
nanocrystals.
Figure 5-6 shows tunneling-conductance spectra measured on two InAs/ZnSe
core/shell nanocrystals with 2 and 6 ML (monolayer) shells, along with a typical
curve for an InAs QD of radius similar to the nominal core radius @ 1.7 nm, deposited on Au with DT linkers. The general appearance of the spectra of core and
core/shell nanocrystals is similar and the band-gap is nearly identical, as observed
in the optical absorption measurements [34]. In contrast, the s-p level separation is
substantially reduced. Both effects are consistent with a model in which the s state
is confined to the InAs core region, while the p level extends to the ZnSe shell. In
this case, the p state is red-shifted upon increasing shell thickness, whereas the s
level does not shift, yielding a closure of the CB s-p gap.
The complementary optical spectroscopy also provides evidence of the reduction
of the s-p spacing upon shell growth, as manifested by the PLE spectra presented
in Figure 5-7. The three spectra, for cores (solid line) and core/shells with 4 ML
and 6 ML shell thickness (dotted and dashed lines, respectively), were measured
using the same detection window (970 nm), corresponding to the excitonic bandgap energy for InAs cores 1.7 nm in radius. The peak labeled III, which as discussed earlier corresponds in the cores to the transition from the VB edge state to
the CB 1Pe state, is red-shifted monotonically upon shell growth. The dependence
of the difference between peak III and the band-gap transition I on shell thickness
is depicted in the inset of Figure 5-7 (circles), along with the 1Se 1Pe level spac-
315
5 Properties
DE (eV)
0.5
PLE Intensity (a.u.)
316
0.4
0.3
2
0
4
6
Shell (ML)
I
0.0
II
0.2
0.4
0.6
0.8
1.0
E-E det
PLE spectra, normalized to peak III,
for InAs cores (solid line) and InAs/ZnSe core/
shell nanocrystals of four (dotted line) and six
(dashed line) ML shells, with the zero of the
energy scale taken at the detection window
(970 nm). The inset depicts the dependence
Fig. 5-7
on shell thickness of the s-p level closure, as
determined by tunneling (open squares) and
PLE (solid circles). The PLE data points
represent the difference between transition III
and transition I (which hardly shifts upon shell
growth).
ing extracted from the tunneling spectra (squares). While the qualitative trend of
red shift is similar for both data sets, there is a quantitative difference, the optical
shift being considerably smaller. This is in contrast to the good correlation between
the optical and tunneling spectra observed for InAs cores, providing an opportunity to examine the intricate differences between these two complementary methods. While the tunneling data directly depict the spacing between the two CB
states, the PLE data in the inset of Figure 5-7 represent the energy difference
between two VB to CB optical transitions. Therefore, evolution of the complex VB
edge states upon shell growth will inevitably affect the PLE spectra. In particular,
a blue shift of the p-like component of the VB edge state upon shell growth will
reduce the net observed PLE shift compared to the tunneling data, consistent with
the experimental observations.
The elegant artificial atom analogy for QDs, borne out by optical and tunneling
spectroscopy, can be tested directly by observing the shapes of the QD electronic
wave functions, as was demonstrated by various techniques for MBE (molecular
beam epitaxy) grown QDs. For colloidal free-standing nanocrystal-QDs, the unique
sensitivity of the STM to the electronic density of states on the nanometer scale
seems to provide an ideal probe of the wave functions. A demonstration of this
317
5.1 Semiconductor Nanoparticles
1
45 (f1)
(b)
p
s
s
p
I (pA)
dI/dV (a.u.)
topography
p
(a)
30
15
0
-1
(c) VB = 0.9 V
0
bias (V )
1
(d) VB = 1.4 V
1.0 (f )
2
2
(e) VB = 1.9 V
I/Imax
-2
0.5
0.0
-5
0
5
X (nm)
(h)
(i)
(j)core
1
shell
"s"
"p"
2
r
(g)
0
0
1
2
r (nm)
Wave function imaging and calculation for an InAs/ZnSe core/shell QD having a
six ML shell. a A tunneling spectrum acquired
for the nanocrystal. b 8 8 nm 2 topographic
image taken at VB ¼ 2:1 V and Is ¼ 0:1 nA.
c–e Current images obtained simultaneously
with the topographic scan at three different
bias values denoted by arrows in a. f1 Crosssections taken along the diagonal of the current
images at 0.9 V (lower curve), 1.4 V (middle
Fig. 5-8
curve) and 1.9 V (upper curve). f2 The same
cross-sections normalized to their maximum
current values. g–j Envelope wave functions
calculated within a ‘‘particle-in-a-sphere’’
model. The radial potential and the energies
of the s and p states are illustrated in the inset
of frame a. g–i Isoprobability surfaces, showing
s 2 (g), px 2 þ py 2 (h), and pz 2 (i). j The square
of the radial parts of the s and p wave functions normalized to their maximum values.
capability is given by recent work on the InAs/ZnSe core/shells discussed above.
Here, the different extent of the CB s and p states, implied by the spectroscopic
results, can be directly verified by using the STM to image the QD atomic-like wave
functions.
Bias-dependent current imaging measurements were performed, as shown in
Figure 5-8 for a core/shell nanocrystal with 6 ML shell. The dI=dV versus V spectrum is shown in Figure 5-8a, and the bias values for tunneling to the s and p states
are indicated. A topographic image was measured at a bias value above the s and
p states, VB ¼ 2:1 V (Figure 5-8b), simultaneously with 3 current images. At each
point along the topography scan the STM feedback circuit was disconnected
momentarily, and the current was measured at three different VB values: 0.9 V,
3
4
318
5 Properties
corresponding to the CB s state (Figure 5-8c); 1.4 V, within the p multiplet (Figure
5-8d); and 1.9 V, above the p multiplet (Figure 5-8e). With this measurement procedure, the topographic and current images were all measured with the same
constant local tip-QD separation. Thus, the main factor determining each current
image is the local (bias-dependent) DOS, reflecting the shape of the QD electronic
wave functions.
Upon comparing the current images, pronounced differences are observed in
the extent and shape of the s and p wave functions. The image corresponding to
the s-like wave function (Figure 5-8c) is localized to the central region of the core/
shell nanocrystal, while the images corresponding to the p-like wave functions extend out to the shell (Figures 5-8d,e), consistent with the model discussed above.
This can also be seen in the cross-sections presented in frame f1 taken along a
common line through the center of each current image, and most clearly in frame
f2 , which shows the current normalized to its maximum value along the same
cuts. Figure 5-8e, taken at a voltage above the p multiplet, manifests a nearly
spherical geometry similar to that of Figure 5-8c for the s state, but has a larger
spatial extent. Figure 5-8d, taken with VB near the middle of the p multiplet, is also
extended but has a truncated top with a small dent in its central region.
An illustrative model aids the interpretation of the current images, assuming a
spherical QD shape, with a radial core/shell potential taken as shown in the inset
of Figure 5-8a [77]. The energy calculated for the s state is lower than the barrier
height at the core-shell interface and has about the same values for core and core/
shell QDs. In contrast, the energy of the p state is above the core-shell barrier and
it red-shifts with shell growth, in qualitative agreement with our spectroscopic
result discussed in the previous section. Isoprobability surfaces for the different
wave functions are presented in Figures 5-8g to 5-8i, with Figure 5-8g showing the
s state, Figure 5-8h the in-plane component of the p wave functions, px 2 þ py 2 ,
which has a torus-like shape, and Figure 5-8i the perpendicular component, pz 2 .
The squares of the radial parts of the s and p wave functions are presented in
Figure 5-8j. The calculated probability density for the s state is spherical in shape
and mostly localized in the core, while the p components extend much further to
the shell, consistent with the bias-dependent current images. Moreover, the different shapes observed in the current images can be assigned to different combinations of the probability density of the p components.
A filled torus shape, similar to the current image of Figure 5-8d taken at the
middle of the p multiplet, can be obtained by a combination with larger weight
of the ( px 2 þ py 2 ) component parallel to the gold substrate, and a smaller contribution of the pz component. The unequal weights reflect preferential tunneling
through the in-plane components. This may result from a tip-induced Stark effect
leading to a small degeneracy lifting, as shown theoretically by Tews and Pfannkuche [78]. A spherical shape for the isoprobability surfaces results from summing
all the p components with equal weights, consistent with the current image measured at a bias above the p manifold (Figure 5-8e). This example of wave function
imaging combined with the tunneling and optical spectra provides visualization of
the atomic-like character of nanocrystal quantum dots.
5.1 Semiconductor Nanoparticles
319
Level Structure of CdSe Quantum Rods
An additional system prepared by colloidal chemistry that is attracting considerable
interest is that of quantum rods that exhibit electronic and optical properties different from QDs. For example, because of their cylindrical symmetry, QRs have
linearly polarized emission, as demonstrated by fluorescence measurements on
single rods [41], leading to polarized lasing [14]. The powerful combination of
optical and tunneling spectroscopy discussed here was also applied to the investigation of the length and diameter dependence of the electronic level structure of
CdSe quantum rods [40].
The central conclusion drawn from the spectroscopic studies, as well as from
calculations employing a multiband effective mass model, is that the QR level
spectrum depends mainly on the rod diameter and not on its length for CdSe
rods with length above @10 nm [40]. This is demonstrated in Figure 5-9 by both
PLE and tunneling data. PLE spectra (left frame) measured on QR samples differing considerably in length but having similar diameter look very similar, whereas
the spectra measured on QRs of different radii appear to be different. This becomes evident by comparing the upper two PLE curves measured on thin rods
(radii @ 1.7 nm) with the lower curves measured on thicker ones (radii @ 3.1 nm).
The tunneling spectra (right frame), acquired in the regime where charging effects
are suppressed, reveal a similar behavior. It can be seen here as well that the band
5.1.1.5
T=10K
II
I
CB1 CB2
4.2 K
b
I
0.4
0.7
1.1
1.4
E-Edet(eV)
Optical and tunneling spectroscopy
of CdSe QRs. PLE spectra for CdSe QRs are
shown in the left hand frame: a 31 1:9
(length times radius) nm, b 11 1:6 nm,
c 60 3:3 nm, d 11 2:9 nm, with the zero
energy representing the position of the
detection window. Relevant optical transitions
are marked. The structure above 0.7 eV is
overlapping peaks of the excitation lamp that
Fig. 5-9
(60x2)
CB3
(30x2.75)
c
d
II III
0.0
(25x2)
dI/dV (arb. u.)
Intensity (arb. units)
VB1
a
(55x2.75)
-2
-1
0
1
Sample Bias (V)
could not be completely normalized out. In the
right hand frame, tunneling spectra for QRs of
different sizes are shown. For clarity, the
spectra were shifted horizontally to align the
CB1 peaks. Between the left- and right-hand
frames, a TEM image of a QR sample, 31 3:9 nm (top, scale bar is 50 nm), and an STM
topographic image of a single QR, 25 2 nm
(bottom), are shown.
2
320
5 Properties
gap depends mainly on, and decreases with, the QR radius, in agreement with the
optical measurements [39].
Optical Gain and Lasing in Semiconductor Nanocrystals
The broad spectral coverage for luminescence from the ultraviolet to the near
infrared offered via tunability of composition, size and shape in semiconductor
nanocrystals presents an obvious advantage for using such materials as tunable
optical gain media and in laser applications. Furthermore, low lasing thresholds
predicted for quantum dots and quantum wires compared with two-dimensionally
confined quantum wells are the basis of present semiconductor diode laser devices
ubiquitous in information and telecommunication technologies [79, 80].
Amplified spontaneous emission was recently observed by Bawendi, Klimov and
coworkers for spherical colloidal CdSe QDs in close-packed films where pumping
with an amplified femtosecond laser source was used to compete with fast nonradiative Auger decay processes [13, 81]. In further studies, films with CdSe nanocrystals were deposited on a distributed feedback grating structure to yield optically
pumped lasing that was tunable through the visible range by changing the nanocrystal size [82, 83].
Lasing was also observed for semiconductor nanocrystals in solution at room
temperature using a cylindrical microcavity, with nanosecond (ns) excitation. A
simple microcavity set-up was used, providing uncomplicated means for screening
colloidal nanocrystals as potential laser chromophores directly in solution [14].
Figure 5-10 presents results of lasing for CdSe/ZnS quantum rods, along with
their structural and optical characteristics. The quantum rods were grown using
the well-developed methods of colloidal nanocrystals. The rod/shell configuration
was chosen, since the growth of ZnS on the organically coated CdSe quantum rods
increases the fluorescence quantum yield from @2% to 14%. A transmission electron microscope (TEM) image of the quantum rods deposited on a grid from hexane solution is presented in Figure 5-10a. Good size and shape monodispersivity is
observed, and the average nanorod length is 25 nm, with an average diameter of 4
nm. The monodispersivity is deemed essential for achieving lasing, by reducing
the inhomogeneous spectral broadening of the gain profile. The optical characteristics of the nanorod sample are summarized in Figure 5-10b.
In Figure 5-10c, emission spectra collected from the cylindrical microcavity,
loaded with a quantum rod sample in hexane, is shown for several laser excitation
intensities, where each spectrum corresponds to a single laser shot. The peak
at 2.33 eV is due to scattering of the pump laser. At low excitation intensities
(e.g. 0.016 mJ), the spectrum resembles the emission spectrum measured for the
quantum rods in solution (Figure 5-10b). At higher intensities, changes in the
spectral shape are observed. Most significantly, above @0.08 mJ a narrow lasing
peak clearly emerges at 1.98 eV, on the red side of the fluorescence peak. The dependence of the intensity of the lasing peak on pump power is shown in the right
inset of Figure 5-10c (full squares) along with the dependence of the fluorescence
peak intensity (empty circles). The lasing peak intensity exhibits an abrupt change
of slope at the onset of laser action, while at the same time, the peak fluorescence
5.1.1.6
5.1 Semiconductor Nanoparticles
b
Absorbance
PL
a
2.0
c
2.8
Intensity (kcounts)
40
40000
Intensity (counts)
2.4
Energy (eV)
pump
30000
detection
30
20
10
0.00
20000
0.05
0.10
0.15
0.20
Excitation Power (mJ)
10000
0
1.8
2.0
2.2
2.4
Energy (eV)
Fig. 5-10 Frame a is a TEM image of the
quantum rod sample (scale bar 50 nm). Frame
b shows the absorption (solid line), and the
emission (dot-dashed line) for the quantum
rods in solution. Frame c shows emission
spectra for the quantum rods loaded in the
microcavity at different pump powers. The
peak at 2.33 eV is the residual scattering of
the pump laser. The pump intensities from
low to high are: 0.016 mJ, 0.038 mJ, 0.075 mJ,
0.088 mJ, 0.10 mJ, 0.15 mJ, and 0.19 mJ. The
narrow peak at 1.98 eV emerging above the
fluorescence is the laser signal that appears at
a threshold pump level of @0.08 mJ. The inset
on the right in frame c shows the intensity of
the lasing peak (filled squares) and the fluorescence peak (empty circles) versus the pump
power. Lasing shows a clear threshold behavior,
as accentuated by the solid lines. The inset on
the left in frame c is a schematic top view of
the cylindrical microcavity experimental set-up,
with the gray shaded area representing the
quantum rod solution occupying the volume
between the fiber and inner capillary walls.
intensity is in effect saturated. This threshold behavior of the lasing peak intensity
provides an additional important signature of lasing [84].
Higher-resolution spectral measurements revealed that lasing occurs in a sequence of discrete modes corresponding to the whispering gallery modes of the
321
322
5 Properties
cylindrical microcavity. Polarized measurements of the lasing for rods versus CdSe
dots in solution showed that, while for QDs the lasing is not polarized, for quantum rods, a linearly polarized laser signal directly related to their symmetry is
detected [14]. These investigations into gain and lasing in semiconductor nanocrystals reveal examples that may lead to further technological developments and
implementation of colloidal nanocrystals in laser and optical amplifier applications.
Acknowledgements
We would like to thank our students and coworkers D. Katz, Y.-W. Cao, S.-H. Kan,
T. Mokari, E. Rothenberg, Y. Ebenstein, M. Kazes, D. Lewis, Y. Levi, A. Sharoni and
T. Wizansky for their important contributions to the work presented in this chapter. The work was supported in parts by the Israel Science Foundation, the USIsrael Binational Science Foundation, and the DIP foundation.
5.1.2
Optical and Thermal Properties of Ib–VI Nanoparticles
Stefanie Dehnen, Andreas Eichhöfer, John F. Corrigan, and
Dieter Fenske
Optical Spectra and Thermal Behavior
The discrete molecular cluster species presented are ideal subjects for the study of
the molecular quantum size effect [85], which is one of the most urgent questions
prompting research in this area of cluster chemistry. One has therefore to explore
the electronic properties of the cluster species, beginning with the interplay of
cluster size and HOMO-LUMO gap. Theoretical investigations gave preliminary
answers to this question [86, 87]. The first dipole- and spin-allowed singlet or triplet excitations of selected ‘‘naked’’ copper selenide clusters [Cu2n Sen ] (up to n ¼ 15)
were calculated on the CIS [88, 89] level (CIS ¼ configuration interaction considering single excitations) for structures previously optimized for the same symmetry
employing the CCSD(T) [90] method (CCSD(T) ¼ coupled cluster approximation
up to single and double excitation with additional consideration of third excitations
by means of perturbation theory) (n ¼ 1; 2) or the MP2 method (MP2 ¼ 2nd order
Møller-Plesset perturbation theory) (n > 2). Triplet excitations were investigated in
order to take into account the fact that relativistic effects (spin-orbit coupling) [91]
are noticeably present in CuaSe compounds. The results are given in Figure 5-11.
The development of the energy values for singlet or triplet excitations with increasing cluster size shows the same trend, the triplet excitation energies being
0.2–0.7 eV lower than the singlet excitation energies. The excitation energies increase with slight oscillation to a value around 6 eV for the hexamer (n ¼ 3) and
decrease afterwards down to a value of ca. 3.8 eV for the [Cu30 Se15 ] species. The
initial increase in the excitation energy values can be explained firstly by a stabilization of selenium p-orbitals by an admixture of s-, p-, or d-contributions of copper
5.1.2.1
5.1 Semiconductor Nanoparticles
E/eV
6
4
2
0
1
5
10
First dipole- and spin-allowed singlet
or triplet excitations of selected ‘‘naked’’ copper
selenide clusters [Cu2n Sen ] (up to n ¼ 15) calculated on the CIS level (CIS ¼ configuration
interaction with single excitations) for structures previously optimized for the same
Fig. 5-11
n
15
symmetry employing the CCSD(T) method
(CCSD(T) ¼ coupled cluster up to single and
double excitation with additional consideration
of third excitations by means of perturbation
theory) (n ¼ 1; 2) or MP2 method (n > 2).
atoms and, secondly, by an increasing completion of the selenium shell around the
copper substructure. The latter reduces the space for electrons in excited states
with s-character. Thus, for [Cu12 Se6 ], significant p-type contributions are observed
for the calculated transitions. The two exceptions (n ¼ 2; 4) of the described trend
feature less compact molecular structures that lead to lower excitation energies.
Having reached a maximum coverage of selenium atoms on the copper framework, the spatial extension of the cluster molecules gains more relevance. This
confirms the assumption of the splitting of the Se-p and Cu-s states to form a
valence and a conduction band. For clusters of ever-increasing size, and finally
bulk Cu2 Se, the first excitation energy should approach a value of ca. 1 eV [92,
93]. According to the quantum chemical investigations, ‘‘naked’’ copper selenide
clusters of medium size should be colorless insulators that gradually adopt semiconductor properties with increasing cluster size. Calculations on small PH3 - or
PMe3 -ligated copper selenide clusters show an increase in the first excitation
energies. Estimating an error of about 1.5 eV, the respective transitions could indeed absorb violet light – in agreement with the red color of small experimentally
observed copper selenide clusters.
The crystals that are obtained from the cluster formation reactions are intensely
colored. The intensity of the color increases when going from sulfur- to seleniumto tellurium-bridged compounds (see below), as expected for an increase in the
covalent or (semi-) metallic binding properties. Small copper sulfide and selenide
clusters form light red, orange or purple crystals. With increasing cluster size, the
color varies from dark red to reddish-black to (finally) black with a metallic sheen.
The optical spectra of some copper selenide cluster compounds were studied by
means of solid-state UV-VIS spectroscopy. The single crystalline samples were
pulverized and embedded in mineral oil between two quartz plates. Figure 5-12
323
324
5 Properties
[Cu26Se13(PEt2Ph)14]
[Cu44Se22(PEt2Ph)18]
[Cu70Se35(PEt2Ph)n]
[Cu140Se70(PR2R')34]
(15)
(23)
(31)
(35)
Fig. 5-12 Solid-state absorption spectra of PEt2 Ph ligated
compounds 15, 23, 31, and 35, which contain 26, 44, 70, or
140 copper atoms. For the measurements, single crystalline
samples, previously pulverized in mineral oil, were placed
between two quartz plates.
shows the absorption spectra of PEt2 Ph-ligated compounds [Cu26 Se13 (PEt2 Ph)14 ]
(15), [Cu44 Se22 (PEt2 Ph)18 ] (23), [Cu70 Se35 (PEt2 Ph)n ] (n ¼ 23; 24) (31) and
[Cu140 Se70 (Pr2 R 0 )34 ] (35), which contain 26, 44, 70, and 140 copper atoms, respectively. The influence of the phosphine ligands should be roughly the same for all
four clusters, and thus the increments in the measured shift of the absorption
band can be ascribed solely to the increasing size and structural changes of the
CuaSe framework.
One observes a shift of the absorption band to higher wavelengths with increasing cluster size, and this agrees in principle with a decrease of the HOMO-LUMO
gap in the same direction. Although the experimental value for the band gap in
Cu2 Se (1.1 eV, i.e. 1127 nm, from optical measurements [93]; 0.37 eV, i.e. 3350 nm,
from electrochemical measurements [94]) is not reached, the tendency can in fact
be explained by the quantum size effect. An approximation of the frontier orbitals
with increasing atomic number is clearly observed. The large width of the cluster
spectra may be caused by the measuring temperature (room temperature) and/or
the charge-transfer contribution to the observed transitions.
It has been known for several years that a CuaP bond cleavage occurs if copper
selenide clusters are heated. The residual black powders thus formed are Cu2 Se
particles that convert to the a-Cu2 Se phase at 600 C. Additionally, XPS experi-
5.1 Semiconductor Nanoparticles
Fig. 5-13
Results of the thermogravimetry measurement of 31 under different pressures.
ments (XPS ¼ X-ray photoelectron spectroscopy) proved that partial removal of
the terminal ligands under the experimental conditions (room temperature, high
vacuum) led to an aggregation of the remaining clusters to larger, still-ligated particles with a very narrow size distribution [95]. This corresponds with quantum
chemical investigations which claimed that the clusters were metastable species
relative to the binary chalcogenide Cu2 E and which computed very low CuaP
binding energies (see above). Thermogravimetric analyses of 35 (Figure 5-13) show
that all phosphine ligands are removed in a single-step process within 20 min. The
required temperature for bond breakage is also dependent on the experimental
pressure.
On carrying out the experiments at ambient pressure in a nitrogen flow, the
ligands dissociate at 226 C. A reduction in the pressure down to 2 106 mbar
causes a decrease in the temperature to 145 C. The observed mass loss value of
33.6% contrasts with the figure of 35.6% for the mass that should theoretically
have left the system. This observation and a 1.2% carbon residue in the formed
copper selenide powders indicate a partial decomposition of the PR2 R 0 ligands. By
following the processes by means of powder diffraction, it was observed that the
intense reflections of the crystalline cluster (2Y in the range 2–10 ) disappeared
after the thermogravimetric experiments were undertaken. In their place, lower
intensity reflections arose in a 2Y range of 10–55 . Sharp reflections of crystalline
a-Cu2 Se were eventually observed after tempering the powder to 600 C. The
Scherrer equation [96] (Eq. (5-1)) provides a relationship between the medial peak
width b and the particle size D.
D¼
K l 57:3
b cos Q
ð5-1Þ
325
326
5 Properties
D: average particle size perpendicular to the reflecting plane
K: atomic form factor (here: 1.0)
l: wavelength of the X-ray radiation (here CuKa : l ¼ 1:5418 Å)
b: peak width measured at half maximum intensity
Q: diffraction angle
Calibrating the peak width at half maximum intensity for the used diffractometer
to LaB6 , one calculates an average particle size of D ¼ 117 Å for Cu2 Se powders
that were generated by heating 35 up to 150 C at 2 106 mbar for 20 min. This
value for D is 7.8 times the size of the [Cu70 Se35 ] core of 35 (15 Å). The results are
another indication of the partial, controlled collapse of the cluster particle lattice.
Optical Spectra in the Solid State
The colors of the microcrystalline powders of various copper telluride clusters
reveal differences in the electronic properties of the compounds. The larger clusters are significantly darker in color than the smaller ones. For a more detailed
investigation, UV-VIS spectra of the stoichiometric clusters [Cu8 Te4 (PPh3 )7 ] (55)
and [Cu16 Te8 (PPhnPr2 )10 ] (57), as well as for the mixed-valence compounds
[Cu16 Te9 (PPh3 )8 ] (73), [Cu23 Te13 (PPh3 )10 ] (76), [Cu26 Te12 (PEt2 Ph)112 ] (79), and
[Cu44 Te23 (PPhnPr2 )15 ] (83) were measured in the solid state [97]. Absorption spectra of the crystalline samples were measured in transmission mode from suspension in Nujol between two quartz plates. As representative examples, the spectra of
57, 73, and 83 are shown in Figure 5-14.
In line with its red color, 57 exhibits charge transfer bands between 600 and 200
nm together with bands that can be assigned to transitions related to the ligands.
The UV-VIS spectrum of 73 also shows similar bands between 600 and 200 nm,
but additionally displays broad and weaker peaks at higher wavelengths, specifically at 670, 750, 1050 and 1350 nm. As outlined above, 73 can be described as a
mixed-valence compound with an overall electron deficiency that is delocalized in
the valence band. In general, compounds containing metal atoms in formally different oxidation states exhibit interesting spectroscopic behavior. Depending on the
degree of delocalization of the charges over the molecular framework, one may
detect additional bands in the near infrared. According to the classification of
Robin and Day [98], which describes the different degrees of electron delocalization between the metal atoms, the copper telluride clusters under investigation
should belong to class IIIa (metal clusters). In this class, metal centers with different valencies cannot be distinguished, and total delocalization of the charge is
achieved. The larger cluster 83 also formally contains an electron deficiency of
two electrons and thus shows broad peaks that are shifted even further to longer
wavelengths than in the case of 73. Since the energy of the mixed-valence transition is strongly correlated with the degree of delocalization, an even better charge
delocalization – i.e. a smaller transition energy – is observed for 83, as might have
been expected for a larger cluster. In 79, the assignment of formal charges resulted
in two Cu 0 atoms and 24 Cuþ centers, which implies a formal electron excess of
two. In principle, this should also cause mixed-valence bands. However, the ab5.1.2.2
5.1 Semiconductor Nanoparticles
2,2
n
1,8
[Cu16Te8(PPh Pr2)10]
[Cu16Te9(PPh3)8]
1,6
[Cu44Te23(PPh Pr2)15]
2,0
absorbance [a.u.]
n
1,4
1,2
1,0
0,8
0,6
0,4
0,2
0,0
200
400
600
800
1000 1200 1400 1600 1800 2000 2200 2400
Wavelength [nm]
The UV-VIS solid-state absorption spectra (mull in
nujol) of copper telluride clusters [Cu16 Te8 (PPhnPr2 )10 ] (57),
[Cu16 Te9 (PPh3 )8 ] (73), and [Cu44 Te23 (PPhnPr2 )15 ] (83).
Fig. 5-14
sorption spectrum of 79 only shows a red shift of the onset of the charge-transfer
bands with respect to 57 from 600 to 800 nm. It cannot yet be established whether
this is general behavior of the two different types of mixed-valence copper telluride
clusters possessing either an electron deficiency or an electron excess.
Finally, one could conclude that in contrast to the clusters containing sulfur or
selenium ligands, copper telluride clusters show a significant tendency to form
nonstoichiometric clusters, with either an electron deficiency or an electron excess.
This is not only evident from the overall charge, but it is also shown by the occurrence of intervalence bands in the optical spectra of these compounds. Structural
considerations as well as calculations on the electronic band structure reveal that a
localization of the different formal oxidations states is not possible. Even though 79
contains a central Cu 6 octahedron with sufficiently short CuaCu distances, the
charge excess cannot be assigned to d 10 –d 10 interactions. The ‘‘additional’’ electron
pair which arises from the formal assignment of charges Cu 1þ and Te 2 is rather
delocalized in an MO which is embedded in the tellurium(5d)–copper(4s) valence
band. The distinct differences from the Cu2 S or Cu2 Se mixed-valence cluster systems can generally be attributed to the small differences between the ionization
potentials (IP), electron affinities (EA) and electronegativities (EN) of copper and
tellurium, which do not allow a discrete separation into Cu 1þ and Te 2 centers
[IP(Cu): 745.4, IP(Te): 869.2; EA(Cu): 118.5, EA(Te): 190.2; EN(Cu, Pauling): 1.9,
EN(Te): 2.1].
327
328
5 Properties
Fig. 5-15 Trend line of miniaturization in microelectronics
according to ‘‘Moore’s Law’’ as a plot of transistor size vs year.
The dotted line implies that hybrid elements for nanoelectronic
devices can overcome the respective size limits. See body text
and the cited references (e.g. [121]) for explanation.
5.2
Electrical Properties of Metal Nanoparticles
Ulrich Simon
5.2.1
Introduction
Modern microelectronics continuously tends toward a higher degree of integration.
Since the fabrication of the first transistor by Shockley, Brattain, and Bardeen
in the late 1940s [99], transistors in microelectronic components have become
smaller and smaller. This is the basis of all the present-day key technologies which
play a role in all areas of our life. Figure 5-15 shows the continuous reduction in
the minimum size of a transistor, known as Moore’s Law [100]. The trend line has
continued since 1960, and it shows that the component size in microelectronic
devices employed in present-day-technology is about 250 nm and thus about 1
million times smaller than the first transistor.
As miniaturization continues, shorter distances between the transistors and related switching elements on a chip lead to increase in the performance speed and
therefore to shorter delay times in the information exchange. The availability of
nanolithographic fabrication techniques makes it possible to scale down to 100 nm
or less [101, 102]. This has already made a great impact on the performance of
traditional semiconductor circuits as well as opening up new possibilities based
upon quantum effects. However, the same has become true for so-called ‘‘metallic’’
electronics also. Now, following the exploitation of charging effects, or so-called
Coulomb effects, in metallic circuits comprising tunnel junctions with sub-micron
sizes, individual charge carriers can be handled. In metals these carriers are elec-
5.2 Electrical Properties of Metal Nanoparticles
trons and, in the case of superconducting materials at low temperatures, they can
be Cooper pairs of electrons. This field has been named Single Electronics (SE)
[103].
Although charging effects were previously observed in granular thin metal films
[104, 105], SE itself was born in the late eighties, when ultra-small metal-insulatormetal sandwich structures (tunnel junctions) and simple systems including these
started to be intensively studied both theoretically and experimentally [106, 107].
The discreteness of the electric charge becomes essential, and the tunneling of
electrons in a system of such junctions can be effected by Coulomb interaction
of electrons, which can be varied by an externally applied voltage or by injected
charges. A profound understanding of this phenomenon, so-called single-electron
tunneling (SET), was developed, however, when the interest in this field was not a
purely academic one. It is an attractive idea to exploit the principles of SE for the
development of logic and memory cells, which in principle could lead to the construction of a computer working on single electrons, realizing a so-called ‘‘singleelectron logic’’ (SEL). As the physical limits of the production technologies in
device fabrication are rapidly approached, it is becoming clear that a continuous
increase in the integration density requires new concepts to achieve component
sizes of tens of nanometers or less, or ideally, the molecular level.
Different approaches to bridge the size gap between conventionally fabricated
circuit elements and truly atomic scale components have been discussed. These
include the utilization of electron transport through single molecules, the concept
of molecular rectification, and the design of atomic relays or molecular shuttles.
The approach that now receives the most attention among the different concepts of
futuristic circuit design is circuitry based on single-electron transistors. This might
justify the strong and visionary claim of Averin and Likharev that ‘‘single-charge
tunneling is the only evident basis for future nanoelectronics’’ [108].
The aim of this Section is on the one hand to acquaint the reader with charge
transport phenomena in nanoparticles, the physical principles of SET, and basic SE
structures where the charging effects occur as intended. On the other hand, it will
be shown by means of selected examples that in further development of the ideas
of SE, Ligand-stabilized nanoparticles of noble metals have been recognized to be
suitable building blocks for single-electron devices [109, 110].
5.2.2
Physical Background and Quantum Size Effect
The availability of small solids like transition metal clusters [111–113] offers a field
of research that raises basic questions: How small must the number of atoms be
for the properties of the original metal to be lost? How does an ordered accumulation of atoms behave when it is no longer under the influence of its ambient bulk
matter? And, perhaps most important in the context of this chapter: What kinds of
applications of these new materials in microelectronic devices can we expect in the
future?
To answer these questions, we may consider the following scenario. If a metal
329
330
5 Properties
particle, initially having bulk properties, is reduced in size down to a few hundreds
or dozens of atoms, the density of states in the valence and conduction bands decreases and the electronic properties change dramatically, i.e., conductivity, collective magnetism, and optical plasmon resonance etc. vanish. The quasi-continuous
density of states is replaced by quantized levels with a size-dependent spacing.
In this situation, the material chosen does not exhibit bulk metallic or semiconducting behavior, since most of the electrons are tightly bound or localized, i.e.,
a small number of delocalized electrons are left to move as metallic electrons over
the volume of the entire nanoparticle [114]. With respect to the size-dependent
decreasing electrical conductivity, Nimtz et al. [115–117] spoke of the so-called
Size-Induced Metal-Insulator Transition (SIMIT) as early as 1986. This refers to a
transition due to geometrical confinement of delocalized states, with a corresponding de Broglie wavelength, when the volume of a metallic particle is strongly
reduced by fractionation, from a starting diameter of 1 mm to an experimental limit
of less than 20 nm, which was quoted in these experiments. This effect was detected by measuring the dielectric absorption in the microwave range of small
particles of solid or liquid indium metal dispersed in a liquid dielectric. Its explanation claims that the boundaries of these particles confine the electrons into
localized states with wavelengths of multiples of half the de Broglie wavelength
l=2, thus leading to the picture of ‘‘particles in a box’’. The corresponding level
spacing was assumed to be smaller than kB T, as the conductivity of these particles
is only slightly affected by temperature. This means that the particles are still metallic, but with a limited number of electrons determined by the particle size. The
first attempt to take advantage of this effect was the design of so-called Cut-Off
Wavelength (COW) transistors [117].
This ‘‘size-quantization’’ effect may be regarded as the onset of the metalnonmetal transition. The size at which the transition occurs depends on the type of
metal and the criterion chosen for metallicity as well as the method of investigation. Remembering that the term ‘‘SIMIT’’ should not be restricted to the change
in the electrical conductivity, but generalized by means of additional indicators at
the very end of metallic behavior, the weakest-bound electrons which are confined
in the quantum box occupy quantized levels analogous to the valence orbitals in
ordinary atoms [114]. Therefore quantum size nanoparticles are also labeled as
‘‘artificial atoms’’ [118–120].
This is the great opportunity for the chemistry of nanoparticles, as the electronic properties are now tunable via the particle size. With respect to the electrical
properties of nanoparticles and of all kind of arrangements built from nanoparticles, the most important property is the amount of energy needed to add one
extra electron to an initially uncharged particle [121–124]. This energy is the socalled ‘‘charging energy’’, it scales roughly with 1=r (r ¼ radius of the nanoparticle), and it has its atomic analog in the electron affinity of a neutral atom.
Since it can be estimated from the charging energy of a macroscopic metallic
sphere scaled down to the nanometer range, a physical effect like Coulomb blocking of tunneling events, which results from the huge charging energy of metal
nanoparticles, is sometimes regarded as a ‘‘classical size effect’’.
5.2 Electrical Properties of Metal Nanoparticles
Based on the concept of charging energy, the following Section describes the
physical principle of single-electron tunneling in terms of classical considerations.
5.2.3
Single-Electron Tunneling
The transfer of electric charges, or in other words electric current, in a metallic
conductor is associated with the motion of a huge amount of free electrons over
the entire conductor. Because of electron-electron and electron-phonon interaction,
the effective charge which passes a cross-section of the conductor per second is a
result of the displacement of individual charges. As a result, in spite of the discrete
nature of the charge carriers, the current flow in a metal is averaged over a huge
amount of charge carriers and is hence quasi-continuous. In contrast, if one deals
with an isolated nanoscale piece of metal, the number of electrons in this is always
integral. The number of electrons can be changed, e.g., by means of direct connection and then disconnection of that nanoparticle to another charged object or a
lead with a different electric potential. The resulting change in the number of
electrons, which is on the macroscopic scale large and uncontrollable, obviously is
also integral. Therefore, the question arises how to transfer electrons in a circuit
which is built up of nanoscale metallic objects one by one in a strictly controlled
manner?
Physically, such a transfer of electrons can be realized by means of quantum
mechanical tunneling. The probability of such a tunneling event depends on the
potential applied to the circuit as well as on the distribution of excess electrons
over the constituting sites. Therefore, an SET circuit should present a number of
reservoirs for free electrons, which should be small and well-conducting islands
separated by poorly permeable tunneling barriers. As long as the size of these socalled mesoscopic islands is larger than the atomic scale, they certainly comprise a
huge amount of free charge carriers. However, handling of individual charges is
still possible if the characteristic electrical capacitance (C) of the island is small
enough. In this case, the presence of one excess electron changes the electric potential DV of the island considerably, i.e., DV ¼ e=C. Then the control of single
tunneling events can be realized in the following scenario. The enhanced potential
could prevent further charging of the island, essentially reducing the probability
of corresponding tunneling, since it would raise the energy of the whole system.
If then the potential of the neighboring islands and/or gates is changed, the prob-
electrode
electrode
island
Fig. 5-16 The tunneling of a single electron between two metal
electrodes through an intermediate island can be blocked if the
electrostatic energy of a single excess electron on the central
island is large compared with the energy of thermal fluctuation.
331
332
5 Properties
ability of tunneling would be restored, and the tunneling of strictly one electron
would then occur. Simply it might be concluded that SE deals with a small and
defined amount of excess electrons on islands changing their distribution over the
islands in time in a desirable way.
In order to realize this practically, the following two principal conditions must be
fulfilled [103, 125, 126]:
First, the insulating barriers separating the conducting islands from each other
as well as from the electrodes should be fairly opaque. The barrier properties
determine the transport of electrons, which obey the rules of quantum mechanics
and therefore are described by wave functions. The coordinate dependence of a
wave function depends on the energy profile of the barrier. If the energy barrier
separating the islands is high and thick enough (e.g., 1 eV and 2 nm, respectively),
it provides an essential decay of the electron wave function outside the island and,
as a result, only a weak overlapping of the wave functions of the neighboring islands occurs within the inter-island space. If, in addition, the number of electronic
states contributing to tunneling is small enough, as is the case for nanoparticles,
then the total exchange of electrons between the islands becomes negligibly small.
This situation is often referred to as the case of small quantum fluctuations of
charge. In spite of a relatively complex rigorous quantum mechanical consideration, quantitatively this condition can illustratively be formulated using the tunneling resistance RT as a characteristic of the tunneling junction. It should be
much higher than the characteristic resistance expressed via fundamental constants, the so-called resistance quantum, R q ¼ h=e 2 ¼ 25:8 kW. Then the electrons
in the island can be considered to be localized, and their number already behaves
classically, although they undergo thermodynamic fluctuations just like any statistical variable.
Second, in order to have these thermal fluctuations small enough and consequently to make the exchange of electrons controllable, the energy associated with
charging by one extra electron should be much greater than the characteristic
thermal energy kB T. This is the already mentioned charging energy EC and it depends on the charge Q , on the size, and on the charge on the capacitances of any
junctions, gates, conductors etc. in the vicinity of the island. As long as the size of
an island exceeds the electrical screening length (a few tenths of a nanometer), the
geometrical dependence of the charging energy is simply incorporated into the
parameter C (the total island capacitance), and the charging energy takes the simple form EC ¼ Q 2 =2C. Then the condition for the Coulomb energy to take precedence over the thermal energy is quantitatively expressed by EC ¼ e 2 =2C g kB T.
Stated simply, the smaller the island the smaller the capacitance and the larger the
EC , as well as the temperature, at which single-electron charging can be observed
experimentally.
When the above-mentioned two conditions are satisfied, the system is characterized by the set of configurations with various distributions of a small integral
number of excess charges over the islands. These configurations, or states, of the
system can be switched by means of tunneling of individual electrons. The characteristic time scale for the switching between the states due to externally induced
5.2 Electrical Properties of Metal Nanoparticles
single-electron transitions depends on temperature and on the energy distribution
over the whole system. This is the so-called recharging time tR , and it ranges from
about tR ¼ 1011 s for typical values of RT ¼ 100 kW and C ¼ 1016 F up to several
hours with extremely large RT in accordance with the relation tR ¼ RT C.
If such an array is connected to a voltage source, the charge Q on the junction
will increase with increasing voltage unless a tunneling event takes place. As long
as jQj < e=2 the junction is in the Coulomb-blockaded state. At jQj b e=2 tunneling is allowed and will happen after a random amount of time because of the stochastic nature of the process. This is reflected in the current-voltage characteristic
of the junction, since the Coulomb blockade appears at voltages below 0.5 e=C and
a corresponding shift of the tunnel current characteristic.
If the array is biased by a constant current I, at jQj b e=2 tunneling will appear
making jQj jump to e=2, and a new charging cycle starts again. This leads to an
oscillation of the voltage across the junction in a sawtooth-like manner with the
fundamental frequency nSET.
nSET ¼ I=e
In this state the tunneling of single electrons is correlated.
This gives rise to the manipulation of charges in a circuit on the single-electron
level and therefore to the employment of this for the creation of, e.g., sensitive
amplifiers and electrometers, switches, current standards, transistors, ultrafast oscillators, or, generally, digital electronic circuits in which the presence or absence of
a single electron at a certain time and place provides the digital information.
5.2.3.1
Working Principles of Metal-Based Single-Electron Tunneling Devices
The Single-Electron Transistor
The simplest circuit which reveals the peculiarities of single-electron tunneling
comprises only one Coulomb island and two leads (electrodes) attached to it. This
geometry presents the double tunnel junction system. If this system works as a
simple on/off switch, a gate electrode may be capacitively coupled to the island.
Such a system has external control of the state of charge, often called the singleelectron transistor. Applying a voltage to the outer electrodes of this circuit may
cause either sequential transfer of electrons into and out of the central island or no
charge transport, i.e., the transistor remains in nonconductive state. The result
depends on the voltage U, applied to the electrodes as well as on the voltage Ug
applied to the gate.
Because of the small total capacitance of the island, C ¼ C1 þ C2 þ C0 þ C0 0
(C1; 2 are the capacitances of the junctions, which for the sake of simplicity we shall
assume to be equal, C1 ¼ C2 ¼ CJ , while C0 , the capacitance with respect to the
gate and C0 0 , the capacitance with respect to other remote conducting objects, will
be assumed to be much less than CJ ) the island has a large Coulomb energy, and
EC g kB T. The total energy change of the system DE while one electron is tunneling in one of the junctions consists of the charging energy of the island itself as
333
334
5 Properties
Fig. 5-17 a The dependence of the timeaveraged current I on U (the IðUÞ characteristic) for a single electron junction.
Coulomb blockade appears at voltages below
e/C. The dotted line indicates ohmic behavior.
b Time development of the junction charge or
voltage imposed by a constant current source.
well as of the work done by the voltage source. For Ug ¼ 0 it is simply given by
DE ¼ Q f 2 =2C Q i 2 =2C qU;
where Q i and Q f are the initial and final charges of the island, respectively, and q
is the charge (not necessarily integral) passed from the voltage source. From the
elementary reasoning it is clear that if, say, Q i ¼ 0 and Q f ¼ e, then q ¼ Ge=2
5.2 Electrical Properties of Metal Nanoparticles
Fig. 5-18 Scanning electron micrograph of a
single-electron transistor with source and drain
(þU/2 and U/2 respectively) feeding the
central island (marked ‘‘Insel’’), which is
capacitively coupled to a gate electrode. The
size of the central island is about 60 nm 60 nm. (This figure is taken with permission
from BMBF report INFO PHYS TECH, No. 22,
1998.)
Fig. 5-19 Circuit equivalent of a simple single- voltage U induces a net flow of charge through
electron transistor consisting of a two-junction the device, the volume of current I being
arrangement and a gate electrode coupled
controlled by the gate voltage Ug .
capacitively to the central island. The transport
(depending on the direction of electron tunneling) and, therefore, DE ¼ e 2 =2C G
eU=2. This means that for e=C the change of energy is positive when e=C a
U a e=C. Hence, electron tunneling would only increase the energy of the system,
and this transition does not occur if the system cannot ‘‘borrow’’ some energy
from its environment, when the temperature is assumed to be low enough.
Therefore there is a Coulombically blocked state of the single-electron transistor
when the voltage U is within the interval given previously and Ug is zero. Outside
this range, the device conducts current by means of sequential tunneling of electrons.
335
336
5 Properties
Fig. 5-20 Idealized periodic rhomb pattern showing the
Coulomb blockade region on the plane of voltages U and Ug .
A transistor conducts only outside the rhombic-shaped regions.
The value of the number n is the number of extra electrons
trapped in the island in the blocked state.
When the gate voltage Ug is finite, the calculation of the energy change due to
one electron tunneling gives another value, because of additional polarization of
the island electrostatically induced by the gate. The result of this consideration is
illustrated in Figure 5-20, where the U vs Ug diagram of the Coulomb blockade is
shown.
The diagram reflects an interesting feature: it is periodic with respect to the
voltage Ug with a period of e=C0 . This results from the fact that every new ‘‘portion’’ of the gate voltage of e=C0 is compensated by one extra electron on the island,
and it returns to the previous conditions for electron tunneling. Therefore at constant bias U, the current through the device is alternately turning on and off with a
sweep of voltage Ug .
The behavior of the transistor outside the Coulomb blockade region also shows
the single-electron peculiarities, especially for the case of a highly asymmetrical
junction, where R1 g R 2 , and C1 g C2 . In Figure 5-21, one can see the IðUÞ characteristics with a step-like structure fading with decrease of U. This so-called Coulomb ‘‘staircase’’ results from the fact that an increase in U increases the number
of channels for tunneling in a step-like manner, allowing an increasingly larger
number of electrons to be present on the island. Another manifestation of charging effects in the SE transistor is the offset of the linear asymptotes by e=C.
The dependence of the transistor current on the gate voltage opens up the opportunity to fabricate a sensitive device which measures either directly an electric
charge on the island or the charge induced on the island by the charges collected to
the gate, i.e. a highly sensitive electrometer. The sensitivity of such an electrometer,
which has already been reached in practice, is of the order of 104 –105 parts
5.2 Electrical Properties of Metal Nanoparticles
The dependence of the time-averaged current I vs U
for an asymmetrical single-electron transistor.
Fig. 5-21
of electronic charge, exceeding by far the charge sensitivity of conventional devices.
And we shall see later that even the association of charges due to adsorption of
molecules on the coulomb island causes the charge transport through such a
single-electron transistor.
Traps and Turnstiles
The idea of an effective control of the tunneling current by means of electrostatic
gating has been developed further in multi-island systems and has opened up new
possibilities for SE. An example of such a circuit is the so-called single-electron
turnstile. In its simplest version, this consists of three islands aligned between two
outer electrodes, as shown in Figure 5-22. It can be considered as two double
junctions connected in series. The very central island of this chain is supplied by
the capacitively connected gate, which controls electron tunneling in all four junctions. However, in contrast to the operation of the single-electron transistor, the
additional two islands allow the realization of this condition. The central island
works as a controllable trap for electrons, endowing the turnstile with its unique
properties.
If the entire chain of islands is initially blocked by a small bias voltage U and the
gate voltage Ug is swept, the central island first attracts one electron from the left
hand side. This becomes possible because the Coulomb blockade in the left double
337
338
5 Properties
Fig. 5-22 Circuit equivalent of a 4-junction (3-island) singleelectron turnstile. The alternating voltage applied to the gate
is Ug ¼ Uamp sinðotÞ, and at its rise an electron makes a
transition. The electron remains trapped in the island until
the negative half-wave of Ug is diminished.
junction is lifted because of the gate voltage applied. The blockade is automatically
restored after this tunneling even when Ug is increased, and the electron thereby
remains trapped in the central island. If then the voltage Ug diminishes the
blockade in the right double junction, the electron sequentially passes two right
junctions, completing its traveling through the whole array. Hence, if one applies
an alternating voltage, e.g., sinusoidal, to the gate of the turnstile, it should transfer exactly one electron during a period of the signal.
Frequency-controlled single-electron tunneling opens up an opportunity to
maintain a d.c. current through a device. Its value is determined by the frequency n
of the alternating field applied to the gate according to the fundamental relation
I ¼ enSET .
This is of great importance for modern quantum metrology, which could then
determine the unit of d.c. current via the unit of the frequency using one universal
constant, i.e. the electronic charge. In that case, taking into account the very high
accuracy of the atomic standard of frequency, the accuracy of this current standard
would be basically limited by the accuracy of the electron transfer cycle.
5.2.4
Single-Electron Tunneling in Nanoparticles
The theoretical background of SET has found its experimental manifestation in
lithographically fabricated capacitors, as found in conventional computer circuits.
These typically have capacitances in the pF range. Because of their extremely low
charging energies, they would need to be cooled down to the sub-mK range for
single-electron operation. Furthermore, such a capacitor is typically driven at an
operating voltage of 10–100 mV, which would lead to the storage of a few tens of
thousands of electrons per charging. Going to advanced electron beam or extreme
UV lithography, which allows fabrication tunnel junctions with a typical size of 50
nm 50 nm or less, devices still have to be cooled down to at least liquid helium
temperature. The utilization of SET events for applications up to the range of room
5.2 Electrical Properties of Metal Nanoparticles
temperature leads to the necessity of decreasing the junction capacitance, at least
by two orders of magnitude. As already predicted almost ten years ago, this can be
realized by the use of ligand-stabilized metal nanoparticles in a size range of a few
nanometers [124]. These nanoparticles are stabilized by organic molecules, the ligands which surround the metal cores and play the role of a insulating layer in
contacts with neighboring clusters as well as with conducting objects. To use these
particles as building blocks for a new nanoscale architecture, generally new techniques for a defined organization of such nanoparticles needed to be developed to
build up single-electron circuits of different complexities. So far, these have utilized the principles of self-assembly by controlling intermolecular interaction, one
of the main interests of supermolecular chemistry, and this will be a key feature in
this development.
In general, the small size of the nanoparticles makes the number of electrons
countable. This requires a slight modification of the two principal conditions for
‘‘classical’’ SET.
First of all, the standard (macroscopic) electrostatic approach, where a conductor
is treated as a continuous entity with infinitely thin screening depth, fails, as the
latter almost reaches the size of the object, and the number of weakly bound or
metallic electrons in the object is small. As a result the electron-electron interaction is not completely screened out, so that the concept of electrical capacitances
cannot strictly be applied. In this case, the calculation of the charging energy
should be done by counting the energy of interacting charges. However, the Coulomb energy can still be roughly described by the same elementary formula for the
charging energy of a capacitor. Then the symbol C denotes a capacitance which
generally depends on the total number of interacting electrons occupying the
nanoparticle.
The second point is that, because of the small size of a nanoparticle, it changes
its energy spectrum essentially because of the previously mentioned quantum size
effect. In contrast to a macroscopic piece of metal, which has a quasi-continuous
electron energy spectrum (as assumed in the discussion of charging effects in
larger structures), a metal nanoparticle presents for electrons a small potential
energy well which in accordance to quantum mechanics has a set of distinctly discrete energy levels. Their density and distribution depend on the size and shape of
the cluster and, consequently, the tunnel junctions. When one of its electrodes is
a nanoparticle, a tunnel junction can hardly be characterized by such a simple
parameter as a constant tunneling resistance RT . Such an approximation is valid
when densities of levels in both electrodes are high and independent of energy.
In nanoparticles of 1–2 nm, the level spacing appears and is on the scale of the
charging energy itself.
These peculiarities of metal nanoparticles do not, however, eliminate charging
effects in them. Theory and experiment show that the fundamental results of
single-electron tunneling still qualitatively hold true, and new features are being
discovered. For example, the current-voltage characteristic of a double barrier
junction with a quantum dot possesses at low temperature, as well as the Coulomb
staircase, a fine structure due to energy quantization inside the dot. Nevertheless,
339
340
5 Properties
the interpretation of these results is complicated by the fact that the characteristic
time RT C becomes as short as the characteristic time of the energy relaxation
inside the dot, and consequently the electrons no longer have the pure Fermi distribution. Thus, the tunneling characteristics of single nanoparticles provide information about their electron energy spectrum, and hence the density of levels, their
degeneracy, etc.
Thin Film Structures and STM Single-Electron Systems
More than a decade ago, fabrication techniques for the realization of the previously
mentioned ‘‘traditional’’ single-electron circuits (like transistors) were developed
[127, 128]. These techniques are based on the evaporation of metal (usually Al)
films through a fine-pattern mask, which is often made of polymer resist, by
means of an electron-beam lithographic process. The mask provides a number of
thin splits, determining the sizes and shapes of resulting wires and islands, and
the bridges interrupting them determine the position of the desired tunnel junctions and gates. Because of the composite polymer resist layer, which is chemically
under-etched, the resulting mask is rigidly fixed above the substrate and its bridges
become suspended above it. Such arrangements of splits and bridges enable metals to be evaporated onto the substrate from various angles, which allows the
resulting evaporated layers to be in different positions.
In order to form a set of islands connected by tunnel junctions, the evaporation
is usually made from two or three different angles in two or three steps, respectively, with an intermediate oxidation process, when aluminum is used. By adjusting the angles, the linear sizes of the resulting overlapping area of the first and the
second metallic layers can be even smaller than the width of the strips, i.e., somewhat less than 100 nm. At the same time, a gate electrode coupled to the islands
only capacitively should not overlap them, and this is taken into account in the
mask pattern design. After double evaporation, the mask and the supporting
layer are lifted off. Figure 5-18 shows such an SE transistor fabricated as an
AlaAl2 O3 aAl stack.
The resulting barrier of Al2 O3 , with a desirable thickness of a few nanometers, is
mechanically and chemically stable. The height of the corresponding energy barrier is about 1 eV, and this readily provides a tunnel resistance in the range of 100
kW for the junctions of the above-mentioned size. The resulting capacitance of
such tunnel junctions is about 1016 F. This means that these circuits operate
reliably only at temperatures below 1 K, which is just attainable with immense
technical effort.
For essential decrease of the tunnel junction size and corresponding increase of
the operating temperature, more sophisticated technologies for fabrication of SE
structures are certainly needed. An alternative route is the use of ligand-stabilized
metal nanoparticles as building blocks. The drastic reduction in the sizes and
therefore in the capacitances of tunnel junctions are basically realized, e.g., in a
double junction system consisting of a small metallic particle lying on a thin
insulating film on a conducting support under the tip of the scanning tunneling
microscope (STM).
5.2.4.1
5.2 Electrical Properties of Metal Nanoparticles
STM tip
U
It
Metal nanoparticle (droplet)
Insulating layer
Metallic ground plate
Fig. 5-23 Single-electron two-junction system consisting of an
STM tip and a metallic nanoparticle as a central electrode on a
ground plane.
One of the junctions, which is between the ground plane and the metal particle,
is mechanically fixed, while another one, which is between the particle and the
STM tip, is adjustable.
The structures described above have been successfully and repeatedly realized in
practice. As an early example, a one-dimensional arrangement forming a doublebarrier SET junction working at room temperature has been realized by an STM
tip above Au particles ca. 4 nm in diameter obtained by metal evaporation on a 1
nm thick layer of ZrO2 (tunnel barrier) on a flat Au substrate [129]. The experimentally determined capacitance of the nanoparticle-substrate junction is about
1018 F. This value is in good agreement with the theoretically estimated value
based on the model of a parallel-plate capacitor with ZrO2 as a defined dielectric.
Dorogi et al. prepared layers of isolated nanoparticles from a multiple expansion cluster source on self-assembled monolayers of xylene-a,a 0 -dithiol (XYL), and
studied the electrical transport through the tunnel junction in an ultrahigh vacuum system with an STM. The tunneling current as a function of applied voltage
yields evidence of SET at room temperature, and by fitting the IðUÞ curves to a
Coulomb blockade model, the electrical resistance of the XYL molecules was estimated to be about 9 MW [130].
This result is in good agreement with most theoretical expectations of 12.5 MW
for these molecules [131]. Based on the elastic-scattering Green’s function theory,
Wang et al. have developed an approach to characterize the electron transport process in molecular devices. This approach allows the description of the coupling
between the molecule (forming the tunnel barrier) and metals (forming the electrodes) at the hybrid density functional theory level. For XYL, IðUÞ curves have
been calculated, and these are found to be in good agreement with experimentally
obtained ones at room temperature.
These examples show that SET in the room temperature range is accessible by
means of metal nanostructures, but, for highly redundant SET devices, particle size
distribution, at least with atomic precision, has to be avoided, which is not possible
when metal evaporation is used as the nanoparticle source.
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5 Properties
STM tip
U
It
Ligand Shell
Pt309-cluster
Substrate (Au (111))
Fig. 5-24 A single ligand-stabilized Pt309 -cluster between an
STM tip and an Au(111) facet. The junction between the cluster
and the substrate is built up by the ligand shell.
SET on Single Chemically Tailored Nanoparticles
The first results obtained on single chemically tailored nanoparticles were reported
by van Kempen et al. in 1995 [132, 133]. They reported scanning tunneling spectroscopy at 4.2 K on a Pt309 phen36 O20 cluster synthesized by Schmid and coworkers.
The IðUÞ characteristics exhibit clear charging effects, which indicate that the
ligands are sufficiently insulating tunnel barriers between the cluster and the substrate. The experimentally observed charging energy varies from 50 to 500 meV,
while a value of 140 meV would be expected assuming a continuous density of
state in the particles and a dielectric constant of er ¼ 10 for the ligand molecules
when the classical formula EC ¼ e 2 =4pe0 er R (where R ¼ radius of the particle) is
applied.
There are different factors, affecting the variation in EC , which have been taken
into account. Since the clusters have different facets (squares and triangles) EC will
depend on the exact way in which the cluster lies on the substrate. Furthermore,
the ligands, which for simplification have been assumed to be in the form of a
spherical coating on the cluster, may have different orientations varying from
cluster to cluster with respect to the underlying substrate, thus causing a different
tunnel barrier between the cluster and the ground, and therefore a different capacitance. Finally, residual water molecules from the solvent may be physically or
chemically bound on the cluster surface or on the ligand shell, and therefore the
tunnel junction between cluster and substrate may differ for different clusters.
Additional structures on the charging characteristic have been observed in some
cases, which might be expected from a discrete electron level spectrum in the
cluster caused by the quantum size effect. The effect of discrete levels on the
charging characteristics has been treated theoretically by Averin and Korotkov
[134], who extended the existing ‘‘orthodox’’ theory of correlated SET in a double
normal-metal tunnel junction to the case of a nanoscaled central electrode. According to this, the IðUÞ characteristics should exhibit small-scale singularities
reflecting the structure of the energy spectrum of the central electrode, i.e. the
5.2.4.2
5.2 Electrical Properties of Metal Nanoparticles
nanoparticle. Furthermore, the energy relaxation rate becomes evident because of
the small recharging time, t, resulting from the small junction capacitance according to t ¼ RT C, where RT is the resistance of the respective tunnel barrier
and C is the capacitance of the junction. While the theoretical prediction for the
level splitting D according to D ¼ 4EF =3N or to a refined approach by Halperin
[135] leads to an assumed splitting of ca. 8 meV (EF is the Fermi energy of the
metal and N is the number of free electrons), the fit of the spectra lies between
20 and 50 meV. Although different reasons are discussed in the cited paper, the
measurements reflect the discreteness of the level spectrum in the clusters as big
as the 2.2 nm Pt309 cluster.
On smaller ligand-stabilized metal clusters of 1.4 nm, i.e. Au55 (PPh3 )12 Cl6 (Au55 ),
scanning tunneling spectroscopy (STS) could be observed up to the room temperature range. Already, cluster pellets, i.e. three-dimensional compacts of the clusters, gave hints of single-electron transfer in the IðUÞ curves [136] taken at room
temperature. But in these investigations neither the vertical nor the lateral arrangement of the clusters is well defined, as would be the case in a one-, two- or
three-dimensional superlattice. As a consequence, the charging energy has a wider
spread, as would be expected from the estimation of the charging energy of a
single cluster according the formula mentioned above.
Chi et al. studied tunneling spectroscopy on Au55 monolayers prepared on various technically relevant substrates [137]. The samples could be obtained utilizing
a two-step self-assembly (SA) process and a combined Langmuir-Blodgett/SAprocess. The spectroscopy gives clear evidence of the Coulomb blockade originating from the double barrier at the ligand-stabilized cluster as the central electrode
up to the room temperature range.
From a fit of the experimental data at 90 K, the capacitance of the cluster was
calculated to be 3:9 1019 F. This value, which is very sensitive toward residual
charges and nearby background charges, is close to the value of the microscopic
capacitance determined earlier by temperature-dependent impedance measurements [109]. Furthermore, these results are found to be in good agreement with
the capacitance data obtained on the above-mentioned gold nanoparticles on an
XYL-modified Au [111] surface [130, 138].
Most recently, low-temperature tunneling spectroscopy on individual Au55 clusters under ultra-high vacuum conditions have been reported [139]. In accordance
with the above-mentioned results, clear evidence of the Coulomb blockade is given.
In this regime the conductivity appears to be largely suppressed, but is not zero.
This is attributed to a certain probability of co-tunneling within the Coulomb gap
at the finite temperature of 7 K. At this temperature, thermal motion is sufficiently
reduced and the molecular structure of the ligand shell is partly visible. The STM
image fits the space-filling model of the cluster fairly well. Both are shown in Figure 5-26.
In this figure the two locations at which the tunneling spectra have been recorded are indicated. One location is just above a C6 H5 ring of the PPh3 ligand and
the other is next to the ring. Figure 5-27 shows the conductivity peaks, which precisely coincide for both spectra.
343
344
5 Properties
Fig. 5-25 SET on a single-ligand-stabilized
Au55 -cluster at 90 K. R1 , C1 and R2 , C2 are the
junction resistance and capacitance of the
junction between the cluster and the substrate
and the junction between the cluster and the
STM tip. HC1 was calculated to be 3 1019 F
by fitting (reproduced with kind permission
from [137]).
Fig. 5-26 A STM image of a single
Au55 [P(C6 H5 )3 ]12 Cl6 cluster on an Au(111)
surface at 7 K, obtained at a bias of 2 V and a
current of 100 pA using a Pt-Ir tip. Image size:
3.3 nm 2.9 nm. B Space-filling model of the
cluster compound : cuboctahedral core with
55 Au atoms (yellow), 12 P(C6 H5 )3 molecules
(with P pink, C grey and H blue) bound to the
12 edges of the cuboctahedron and 6 Cl atoms
(violet) located in the center of the 6 square
faces of the Au core. Comparison between (a)
and (b) shows that the STM resolves the C6 H5
rings. Spectroscopic data were acquired at the
two distinct locations marked in A. (Printed
with permission from [139].)
5.2 Electrical Properties of Metal Nanoparticles
Fig. 5-27 Tunneling spectra acquired at the
two distinct locations, a and b, marked in
Figure 5-26A. The dashed curve was taken just
above a C6 H5 ring and the solid one next to
the ring. The bias refers to the substrate
potential. The arrows indicate conductivity
peaks, which precisely coincide for both
spectra. (Printed with permission from ref.
[139].)
This shows that the discrete energy levels of the cluster become visible in terms
of conductivity oscillations with an average level spacing of 135 meV. If this value
is compared to that expected for a simple free-electron model, as described above,
the ‘‘electronically apparent’’ cluster diameter is ca. 1.0 nm. This is significantly
smaller than the geometrically determined diameter of 1.4 nm for Au55 and
slightly larger than the expected value of 0.84 nm for Au13 . The author claims that
the difference is most likely caused by the six Cl atoms located at the six square
faces of the cuboctahedral Au55 surface. Since Cl has a high electronegativity it removes one electron from the Au55 core. Nevertheless, within the discrete energy
levels the cluster exhibits metallic behavior.
Metal Nanoparticles in Single-Electron Devices
In order to utilize the opportunities which arise from the room-temperature
SET peculiarities for applications beyond laboratory experimentation, two different
routes have principally been followed with respect to the experimental design. On
the one hand, techniques of fabricating metallic electrodes have been developed
which enable the addressing of a few or even single nanoparticles by electron beam
lithography. On the other hand, STS techniques have been developed further for
chemical control of gating of the central island in an SET transistor or for the
chemical switching of the tunneling barrier, i.e. the insulating organic molecules,
separating the nanoparticles from the underlying support.
5.2.4.3
345
346
5 Properties
5.2 Electrical Properties of Metal Nanoparticles
Mesoscopic Arrays with Nanoelectrodes
Alivisatos and coworkers realized an electrode structure scaled down to the level of
a single Au nanoparticle [140]. To fabricate this structure, they combined optical
lithography and angle evaporation techniques (see previous discussion of SET device fabrication). A narrow gap of a few nanometers between two Au leads on an Si
substrate was defined. The Au leads were functionalized with hexane-1,6-dithiol,
which binds linearly to the Au surface. The free end, which faces the solution,
was used to immobilize 5.8 nm Au nanoparticles from solution between the leads.
The resulting device reveals slight current steps in the IðUÞ characteristic at 77 K
(see Figure 5-28). By means of curve fitting to classical Coulomb blockade models,
the junction capacitances were found to be 2:1 1018 F, respectively, and the
resistances 32 MW and 2 GW, respectively. In a further development of this device,
a gate electrode to externally control the current flow through the central island, a
CdSe nanoparticle, was applied to realize an SET transistor [141].
Sato et al. [142] reported detailed electrical characteristics an SET transistor
utilizing charging effects on metal nanoparticles. This was fabricated by means of
metal electrodes, formed by electron beam lithography, to which a self-assembled
chain of colloidal gold nanoparticles was connected. The inter-particle connection
as well as the connection to the electrodes results from a linkage by bifunctional
organic molecules, which provide the tunnel barriers.
In this device, gold nanoparticles with an average diameter of 10 nm were
deposited on a thermally grown SiO2 surface on an Si substrate by using alkanesiloxane molecules as an adhesion agent. After formation of the SiaOaSi bond
by thermal treatment, terminal amino groups of the silane became attached to gold
nanoparticles to form a sub-monolayer. After this coating, where a sub-monolayer
was formed, 1,6-hexanedithiol was added to interconnect the particles with a more
defined spacing. With a second immersion in a gold particle solution, the coverage
was complete, and the dithiol molecules maintained the distance between the particles, while the second layer filled the gaps between the particles of the first, where
however, chains of 2–4 particles were formed. When this procedure was performed
on an SiO2 substrate equipped with source, drain and gate metal electrodes defined by electron beam lithography, the particles formed a chain of at least three
particles bridging the gap between the outer driving electrodes.
5.2.4.4
Fig. 5-28 a Field emission scanning electron
micrograph of a lead structure before the
nanocrystals are introduced. The light gray
region is formed by the angle evaporation, and
is ca. 10 nm thick. The darker region is from a
normal angle evaporation and is ca. 70 nm
thick. b Schematic cross-section of nanocrystals bound via a bifunctional linker
molecule to the leads. Transport between the
leads occurs through the mottled nanocrystal
bridging the gap. c I–Vsd characteristic of a
5.8 nm diameter Au nanocrystal measured at
4.2 K. The solid lines show three I–V curves
measured over the course of several days. Each
is offset for clarity. These different curves result
from changes in the local charge distribution
about the dot. The dashed lines are fits to the
data using the orthodox Coulomb blockade
model as discussed. d I–Vsd characteristic of
a 5.8 nm diameter Au nanocrystal measured
at 77 K. Several Coulomb steps of period
DVsd @ 200 mV can be seen. (Printed with
permission from [140].)
347
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5 Properties
Fig. 5-29 SET transistor, based on self-assembling of gold
nanoparticles on electrodes fabricated by electron beam
lithography. (Reproduced in modified form from [142].)
Since not all the steps of this procedure could be controlled in detail, the number
of nanoparticles in the bridge chain differed from device to device, but in any case
electron conduction dominated by single-electron charging indicated by a Coulomb
gap could be observed up to 77 K, whereas the nonlinearity is smeared out at room
temperature. Simulation and fitting of the data with a model circuit for a threedot (four junction) SET-transistor indicated that the capacitance of all junctions
in the chain was 1.8–2 1018 F, and the calculated Coulomb gap was in rea-
5.2 Electrical Properties of Metal Nanoparticles
Fig. 5-30 Schematic representation of the
set-up for single particle measurements. Pt
denotes two free-standing Pt electrodes
(dashed region). A ligand-stabilized Pd cluster
is polarized by the applied voltage and
attracted to the gap between the Pt electrodes
(electrostatic trapping: ET). (Printed with kind
permission from [144].)
sonably good agreement with the value of 150 mV obtained from the measured
IðUÞ characteristic, which was systematically squeezed when a gate voltage of
0.4 to 0.4 V was applied. The plot of the current through the device was
clearly dependent on the gate voltage, showing the typical current oscillations,
proving that the desired function of the single-electron transistor had been achieved. According to a Green’s function-based method, which was proposed by
Samanta et al. [143], the transmission function of electrons across the dithiol
ligands was calculated, from which the resistance R per molecule was calculated
by the Landauer formula, R ¼ ðh=2e 2 Þ=TðEg Þ, where TðEF Þ is the transmission
function, T A exp½2ðmEg Þ 1=2 =h. Assuming a barrier height of Eg A 2:8 eV in the
dithiol molecules, the resulting resistance was estimated to be R A 30 GW.
Schmid and Dekker reported a new technique, which allows the controlled
deposition of a single nanoparticle between two metal nanoelectrodes, i.e. the
technique of electrostatic trapping (ET) [144]. The method is based on the attraction of a polarized metal nanoparticle to the strongest point of the electric field,
which was applied to two Pt electrodes.
From solution, the particles can be immobilized in the gap between the Pt electrodes. The latter can be reduced down to 4 nm. The fabrication of the electrodes
was by thermal growth on Si. A 60 nm SiN film was deposited on top, and electron
beam lithography with PMMA and reactive ion etching was used to open a 100 nm
slit in the SiN film with a local constriction with 20 nm spacing. Under-etching
with HF enables free-standing SiN ‘‘fingers’’ to form, which were sputtered with Pt
to reduce the gap down to 4 nm.
In this gap, Pd nanoparticles were trapped to study the electrical transport properties of this double-barrier system. A typical IðUÞ curve is show in Figure 5-32. At
4.2 K, the most pronounced feature is the Coulomb gap at a voltage of about 55 mV,
which disappears at 295 K. Furthermore, the IðUÞ curve is not linear above the gap
349
350
5 Properties
Fig. 5-31 a Pt electrodes (white) separated by a ca. 14 nm gap.
b After ET, the same electrodes are bridged by a single ca. 17 nm
Pd particle (printed with kind permission from [144]).
Fig. 5-32 Current-voltage curves measured at
4.2 K (open squares) and 295 K (solid squares).
The solid curves denote fits of the KN model.
Fitting parameters for these curves are Vc ¼
55 mV, R 0 ¼ 1:1 10 11 W, q0 ¼ 0:15e (offset
charge), and a ¼ E c/h ¼ 0:5. The dashed
curve (a ¼ 0) represents the conventional
model which assumes a voltage-independent
tunnel barrier (printed with kind permission
from [144]).
5.2 Electrical Properties of Metal Nanoparticles
voltage, but increases exponentially, which was explained by suppression of the
effective tunnel barrier by the applied voltage.
To increase the temperature at which single-electron charging in nanoelectrode
arrays may occur, smaller ligand-stabilized Au nanoparticles (Au55 (PPh3 )12 Cl6 )
were deposited by ET in a three-electrode arrangement [145]. The device consisted
of three tungsten electrodes, fabricated from a 25 nm tungsten layer on an Si wafer
with an 80 nm thick SiO2 layer using an electron beam lithography process. The
minimum component size turned out to be 50 nm, with gaps between the contacts
of about 20–50 nm. Upon ET, quasi-one-dimensional chains of the 1.4 nm Au
particles are formed. These chains exhibit Coulomb blockade effects between
0.5 V and þ0.5 V at room temperature, but without any singularities in the
characteristic, which could refer to the electronic structure of the particles.
a)
contact pad
tungsten nanocontact
b)
Fig. 5-33 Complete nanostructure with contact pads (the inset
is nanostructured with three nano tips). SEM image of the
tungsten nanocontact (b) Au55 cluster deposition (printed with
kind permission from [145]).
351
352
5 Properties
Chemical Switching and Gating of Current through Nanoparticles
A significant advantage of incorporating nanoparticles in single-electron devices
is the opportunity to chemically control the size of the particles as well as the
thickness, composition, and state of charge of the ligand shell. The current flow
through such a device will thus be very sensitive to any charges and impurities that
reside on the nanoparticles or in the ligand shell. This sensitivity can be used to
switch the ‘‘transparency’’ of the ligand shell or to use the ligands as ‘‘chemical
gates’’ to manipulate the SET current.
Most recently, Schiffrin and coworkers have described how to control the transparency of the insulation barriers between a substrate and a nanoparticle in an STS
experiment [146].
A bipyridyl moiety (viologen group, V 2þ ) was used as a redox group incorporated
in the ligand shell of the particle. Electrons were incorporated into this group
under electrochemical control, while the transparency of the insulating barrier
was measured by STS. They found that reduction of V 2þ to the radical V.þ leads
to a significant decrease in the barrier height. Further reduction to V o results in a
very large increase in the barrier height. This result reflects the supporting effect
of a half-filled molecular ortbital in V.þ . This might lead to an extension of the
electronic wave function from the nanoparticle to the substrate via the orbitals of
the radical V.þ . As soon as V o is formed, electron pairing in the LUMO (lowest
unoccupied molecular orbital) suppresses the direct electronic interaction.
5.2.4.5
STM tip
Au nanoparticle
e
-
Redox gate
}
... S
8
N
+
N
2 Br
Au substrate
Fig. 5-34 Scheme describing the redox switch, which is based
on a viologen redox center incorporated within the nanoparticles ligand shell. For simplification the counterelectrode
is not shown (adapted from [146]).
-
S ...
+
8
5.2 Electrical Properties of Metal Nanoparticles
These results show impressively that switching of the SET current through a
ligand-stabilized nanoparticle can be induced by electron injection into a specific
redox group within the barriers of the tunnel junction. While the configuration,
which was studied in this work, requires reduction by at least 30 electrons to
change the transparency of the barrier in the nanoparticle layer, a great challenge
will be to integrate these switching elements into a self-assembled SET circuit.
Feldheim et al. have been considering the possibility of employing particlecapping ligands as ‘‘chemical gates’’ to control the SET current flow in an STS
configuration [147]. They reported the IðUÞ characteristics of ligand-stabilized Au
nanoparticles in aqueous solutions, whereas octanethiol-stabilized 5 nm particles
and galvinol-stabilized 3 nm particles were used at the central island. Galvinol
represents a pH-active probe with pK a A 12, whereas about 15 galvinol ligands
were introduced into the ligand shell. Upon increasing the pH, galvinol, however,
is converted to the galvinoxide anion, making the ligand shell negatively charged.
The authors present a collection of IðUÞ characteristics for galvin and octanethiol
nanoparticles in H2 O with several clear current steps and voltage plateaus, reflecting SET in the individual particles. While only small differences in the positions
and magnitudes of the staircase voltage plateaus appear for octanethiol-stabilized
Au upon changing the pH, galvin reacts sensitively when the pH changes from 5
to 8, 10, and finally 12. The authors point out that, firstly, a subtle shift (ca. 30 mV)
in the entire staircase to positive bias potentials is noticeable from pH 5 to pH 8.
The shift is even more pronounced in IðUÞ-curves obtained at pHs 10 and 12
(from ca. 60 to 120 mV). Secondly, DV ¼ e=c decreases in magnitude with increasing pH. Thirdly, slight peaks are evident in the IðUÞ curves at pH b 8. These
chemically induced changes, which lead to an increase in the negative charge on
the cluster, have two consequences: (i) the total self-capacitance of the nanoparticles decreases from ca. 2:2 1018 F at pH 5 to ca. 0:31 1018 F at pH 12, and
(ii) the induced negative charge, by forming galvanoxide, acts like a negatively
charged gate in solid-state SET devices, causing a shift of the Coulomb blockade to
positive bias potentials. The authors conclude that the fact that potential shifts and
capacitance changes are only observed for galvin-stabilized Au and not for octane
thiol-stabilized Au supports the notion of pH-gated SET.
These results have shown for the first time that SET is dependent on the nature
of the ligands stabilizing the nanoparticles as well as on the composition of the
solution surrounding the function array. This will be useful in studying chemical
signal transduction, where SET currents should be sensitive to single redox or
analysis events.
5.2.5
From Single-Particle Properties to Collective Charge Transport
The single-particle properties discussed in the previous section showed the significance of the high charging energy of metal nanoparticles as a prerequisite for SET
events at elevated temperatures. Together with the limited number of free electrons, this may lead us to regard them as artificial atoms. This raises fundamental
353
354
5 Properties
Fig. 5-35 Current-voltage curves for C8 aAu
(left) and GalaAu (right) in H2 O as a function
of pH (adjusted with phosphate buffer, see
[104]). The numbers 1–4 in the GalaAu data
identify voltage plateaus. Cartoons of the
experimental arrangements for measuring I–V
curves of individual nanoclusters in solution
are shown at the top of each data column. The
insulated STM tip, ligand-capped Au nanocluster, an octanethiol-coated planar Au substrate are shown. Length and shapes are not to
scale (printed with kind permission from [147]).
5.2 Electrical Properties of Metal Nanoparticles
questions about the design of ‘‘artificial molecules’’ or ‘‘artificial solids’’ built up
from these nanoscale subunits [148–151]. Remacle and Levine reviewed these
ideas associated with the use of chemically fabricated quantum dots as building
blocks for a new state of matter [114].
Students in their first chemistry course learn that the simplest molecule H2 is
formed by the overlap of the electron wave functions centered on the individual
hydrogen atoms. Correspondingly, the wave functions of artificial atoms can also
overlap, and electrons can coherently and reversibly exchange between them, which
is the basis of a covalent bond. Extending this idea further, the ordered assembly of
identical nanoparticles in one, two, or three dimensions represents the formation
of an artificial solid or superlattice. The fabrication of two-dimensional and threedimensional ordered superlattices has been demonstrated in several publications
[152–160]. Furthermore it is possible to mix nanoparticles of different chemical
compositions, of bimodal size or of different shapes to obtain tailored nanoalloys
[161]. Levine called these new products ‘‘designer materials’’ with tunable properties [114]. Such artificial solids exhibit delocalized electron states, depending on
the strength of the electronic coupling between the adjacent nanoparticles. The
latter will depend on the size of the nanoparticles, the nature and the covering
density of the organic ligands, the particle spacing, and the symmetry of packing.
As it is known from solid state physics, any kind of disorder affects the electronic
structure of a solid. Accordingly, size distribution, packing defects, and chemical
impurities lead to modification of the electronic structure of the quantum dot superlattice. And even in the case where the size of the nanoparticles is identical over
microdomains of the assembly, slight deviations in the ligand density or orientation mean that the nanoparticles are not strictly the same at each lattice site, as
would be the case for atoms. Therefore, solids consisting of chemically tailored
nanoparticles are inherently disordered, and the loss of translational symmetry affects the extended states in the solid, which is immediatedly reflected in the optical
and electrical properties. Remacle et al. illustrated the crucial point of symmetry
breaking by comparing the experimental and theoretical second-order harmonic
generation (SHG) response of a two-dimensional, hexagonal array of Ag particles
(R ¼ 1:5 nm) with variable interparticle spacing D where D=2R is in the range
1.05–1.5 nm. At large interdot spacing (D=2R ¼ 1:5), the wave function will be
localized on the indivial nanoparticles. If the 2D layer is compressed, the wave
function is delocalized over small domains. A quantum phase transition to a fully
delocalized state is achieved for closely packed nanoparticles at D=2R ¼ 1:05 (see
Figure 5-36).
A corresponding response has been obtained from complex impedance measurements on a monolayer of Ag nanoparticles with 2R ¼ 3:5 nm.
Heath and coworkers, analyzing the complex dielectric modulus, which is the
inverse of the complex dielectric permittivity (for details see [163]), reported a reversible metal-insulator transition. They observed a transition from RC relaxational
behavior to an inductive metallic-like response, when the monolayer is compressed
to decrease the interparticle spacing to <0.6 nm. This suggests that a sequential
transition from hopping transport between localized states via tunneling to metallic transport appears, accompanied by a rapid decay in the relaxation time, when
355
356
5 Properties
Fig. 5-36 Transition from a localized to a
delocalized state in a two-dimensional
hexagonal array of ligand-stabilized 3 nm Ag
particles upon mechanical compression. The
transition is probed by the second-harmonic
response, which is sensitive to the extent of
delocalization of the electronic wave function
(printed with kind permission from [114]).
the layer is compressed. The latter may result from an increasing tunneling rate,
i.e., the resistance decreases between the particles as well as in extended domains
with delocalized states, and this is essentially affected by the degree of disorder.
In order to discuss the signatures of localization and delocalization and its significance to the application of nanoparticles in microelectronic devices, the following Sections give examples of the electrical properties of nanoparticle arrays, distinguished according to their dimensionality.
One-Dimensional Arrangements
In the previous Sections, two examples of one-dimensional (1D) arrays of nanoparticles were described, where self-assembly or electrostatic trapping were used to
address the arrays for electrical transport measurements. A priori it is clear that
both methods cannot lead to strictly 1D defect-free arrangements. Furthermore,
inherent disorder cannot be avoided.
To analyze the effect of disorder, an exact analytical solution in terms of Green’s
function (GF) for the potential distribution in a finite 1D array turned out to be
appropriate [164, 165]. The GF approach allows the formulation the so-called
partial ‘‘solitary’’ problem of small mesoscopic tunnel junctions, similarly to the
problem of the behavior of an electron in 1D tight binding and in a set of random
delta-function models.
The discussion of 1D metal cluster arrangements of ligand-stabilized nanoparticles, as reported in [164, 165], assumes that (i) either the capacitance C is the
same for all junctions, whereas the self-capacitance C0 can fluctuate from site to
site because of a finite size distribution or (ii) the C0 is the same, whereas C can
5.2.5.1
5.2 Electrical Properties of Metal Nanoparticles
Fig. 5-37 a Complex impedance response
(plot of the imaginary part of the complex
modulus M 00 vs. the real part M 0 in the
complex plane) of a monolayer, 3.5 nm in
diameter, of propanethiol-capped Ag nanoparticles. The particle film response is characterized initially by an RC circuit equivalent, in
conformity with a picture of capacitively determined hopping localized conductivity. As the
particles are compressed to a separation of
less than 0.6 nm, the film becomes inductive,
indicating the presence of metallic-like
transport in the film. b The Langmuir isotherm
of the monolayer compression with labels
corresponding to the various curves in the
top graph. Although the isotherm indicates a
collapse near the 130 cm 2 film area, the
collapse is heterogeneous, occurring only at
the mobile barriers. The central region of the
film, where transport measurements were
carried out, remains an uncollapsed monolayer
(redrawn and modified from [162]).
357
358
5 Properties
Fig. 5-38 a 1D array of N ligand-stabilized metal clusters and
b the corresponding circuit equivalent (printed with kind
permission from [164]).
fluctuate from site to site because of packing defects. It has further been assumed
that the metal nanoparticles have a continuous density of states, i.e., quantum size
effects and their influence on the capacitance have been excluded from consideration. These simplifications illustrate the practical implications of disorder with
respect to future applications in microelectronics, but generally they are not necessary assumptions for the method discussed.
The main results are that a decrease in the particle size at one position of the
array increases the potential at this point, which may lead, at least, to localization,
i.e., the single excess electron in the array may be trapped. As a packing defect,
which affects the inter-particle capacitance at one point, acts like an inhomogeneity, the soliton will interact with its mirror-image soliton (or anti-soliton) and
will therefore be attracted. This method is of practical use because the total reflection amplitude obtained by these calculations is directly related to the Landauer
resistance [166–168], which increases exponentially with the number of defects.
For illustration, let us consider a one-dimensional chain of 2 nm neutral particles.
In this chain, a number N of the nanoparticles is replaced by smaller ones (1 nm).
The 1 nm particles represent defects, which lead to a decreased transmission
probability for electrons in the periodic chain, leading to an increase in the overall
resistance. This situation is reflected in Figure 5-39, where the increase in resistance DG1 is given for this selected example, according to [169].
Two-Dimensional Arrangements
Janes et al. reported the electronic conduction across a network of 4 nm gold nanoparticles interconnected by di-isonitrile ligands (1,4-di (4-isocyanophenylethynyl)-2ethylbenzene) to produce a conjugated, rigid molecule with an approximate interparticle spacing of 2.2 nm [170]. The two-dimensional assembly results from the
5.2.5.2
5.2 Electrical Properties of Metal Nanoparticles
359
DG-1 = 2.1 x 105 Ohm
DG-1 = 5.2 x 104 Ohm
DG-1 = 1.3 x 104 Ohm
N=1
N=2
N=3
One-dimensional chain of 2 nm metal particles with
N ¼ 3 ‘‘impurities’’, i.e. 1 nm metal particles. The arrows
indicate that the resistance increases exponentially with N.
Fig. 5-39
deposition of the nanoparticles from a colloidal solution, followed by reactive addition of the linking molecules. In order to allow electrical characterization, the
layers were deposited on an SiO2 -supported GaAs wafer with gold contacts with
separations of 500 nm and 450 nm, respectively.
The IðUÞ-characteristic of the arrays is linear over a broad voltage range. Since
a dot-to-dot capacitance of 2 1019 F may be assumed, the total dot capacitance
will be 1:2 1018 F, if each cluster is assumed to have six nearest neighbors. The
corresponding charging energy will thus be approximately 11 meV, which is only
about half of the characteristic thermal energy at room temperature. Therefore a
Coulomb gap at room temperature is not apparent.
On the other hand, the charging energy of the nanoparticles has been obtained
from comparable samples with two-dimensional cluster linkage by Andres et al.
[171] from temperature dependent d.c. measurements. According to the Arrhenius
relation G0 ¼ Gy expðEA =kB TÞ, where Gy is the conductance as T ! y, EA is
the activation energy and kB is the Boltzmann’s constant, Coulomb charging behavior with a charging energy (which corresponds to the activation energy) of
EA ¼ 97 meV was indicated. The interparticle resistance was revealed to be
0.9 MW, from which a single-molecule resistance of 29 MW is obtained. This value
is in good agreement with the prediction of 43 MW obtained from Hückel-MO
calculations.
Three-Dimensional Arrangements
Signatures of a transition from single-particle properties to three-dimensional
charge transport have been reported by Pileni and coworkers [172]. They reported a
tunneling spectroscopy study at room temperature of 4.3 nm Ag nanoparticles
stabilized in reverse micelles. The electrical transport measurements though single
particles deposited on Au (111) substrates showed that, on increasing the applied
voltage, the small capacitances of the double junction are charged and the detected
current is close to zero because of the Coulomb blockade effect. The IðUÞ curve as
well as the dI=dU vs U plot (insert in Figure 5-40B) clearly indicate the Coulomb
blocking state, i.e., a zero bias and the onset of conductivity through the nanoparticles above the threshold voltage. When the particles are deposited from solutions at higher concentrations, they arrange in a hexagonally ordered monolayer.
The tunneling current though this layer shows a less pronounced nonlinearity in
5.2.5.3
360
5 Properties
Fig. 5-40 Constant-current mode STM
images of silver nanoparticles, isolated (A),
self-organized in close-packed hexagonal
network (C), and in fcc structure (E), deposited
on Au(111) substrate (scan size: A 17:1 17.1 nm 2 , Vt ¼ 1 V, It ¼ 1 nA, C 136 136 nm 2 , Vt ¼ 2:5 V, It ¼ 0:8 nA,
E 143 143 nm 2 , Vt ¼ 2:2 V, It ¼ 0:72 nA).
The graphs are I–V curves and their derivatives
in the inserts of the silver nanoparticles,
isolated (B), self-organized in close-packed
hexagonal network (D), and in fcc structure
(F), deposited on Au(111) substrate [172].
5.2 Electrical Properties of Metal Nanoparticles
the IðUÞ curve as well as a smaller Coulomb gap, as compared to the isolated particles. This indicates the onset of ohmic contributions to the overall conductance of
the system. Going to a three-dimensional arrangement of the particles, which form
a self-organized closed-packed fcc structure, the Coulomb gap almost vanishes, and
the IðUÞ characteristic is a purely ohmic one at low voltages. This reflects the fact
that in the 3D systems an additional transport path for electrons appears between
the adjacent nanoparticles, so that the Coulomb gap between the individual particles is smeared out because of the coupling of the electron states in the closepacked system.
The electrical behavior of the three-dimensional system is also reflected in the
electrical d.c. and a.c. response of compacts of ligand-stabilized nanoparticles [109].
As a common feature, at high temperatures, i.e., several tens of Kelvin below room
temperature, the temperature-dependent d.c. and a.c. conductivities follow a simply activated behavior according to the Arrhenius relation.
sðTÞ ¼ so þ exp
EA
;
kB T
where EA is the activation energy. EA becomes temperature dependent when the
measuring temperature is decreased. This means that down to very low temperatures the conductivity follows the Variable Range Hopping (VRH) expression
proposed by Mott [173]:
To g
sðTÞ ¼ so þ exp
T
whereas g ¼ 1ðd þ 1Þ in d dimensions. Although g ¼ 14 ðd ¼ 3Þ might be expected
from this general expression, g ¼ 12 is predominantly observed in the case of the
compacted metal cluster compounds. This temperature dependence, as well as the
electric field-dependent conductivity, which reflects pronounced non-ohmic behavior at strong electric fields, reveal a pronounced similarity to various heterogeneous
materials, like cermets, doped and amorphous semiconductors, or metal- and
carbon-insulator composites. This behavior was carefully analyzed by van Staveren
and Adriaanse [174–176], applying different physical models of hopping conductivity, with the conclusion that the experimental data can best be fitted by a thermally activated stochastic multiple-site hopping process, whereas at high temperatures around room temperature nearest neighbor hops predominate.
This means that at low temperature the number of charge carriers participating
in the hopping process does not change with temperature. Instead, at high temperatures, where kB T becomes comparable to the charging energy Ec of the metal
particle (note that Ec is determined by the total capacitance of the particle, and is
therefore dependent on the particle size as well as on the inter-particle spacing, i.e.
the dot-to-dot distance), thermally excited extra charge carriers are participating in
the hopping process. Thus, even at high temperatures, the activation energy re-
361
362
5 Properties
flects the energy needed to transfer one electron from one electrically neutral particle to another. This is the charge disproportionation energy, and it has its atomic
analog in the disproportionation energy of atoms or molecules in chemical redox
reactions. At low temperatures, the hopping transport would be expected to become zero if all particles become electrically neutral. Since all reports on the electrical conductivity reflect a residual conductivity at very low temperatures, again it
becomes obvious that packing, shape and size distribution prevent the localization
sites from being identical, as has already been pointed out at the beginning of this
Section.
In terms of a modified Anderson-Hubbard Hamiltonian [177, 178], orientational
disorder and packing irregularities will lead to a distribution of the on-site Coulomb interaction and of the interaction of electrons on different (at least neighboring) sites. This is explicitly pointed out by Cuevas, Ortuno and Ruiz [179]. In contrast to the Coulomb-gap model of Efros and Sklovskii [180], they took into account
three different states of charge of the mesoscopic particles, i.e. neutral, positively
charged, and negatively charged. With this model, the VRH behavior, which dominates the electrical properties at low temperatures, can be fully explained.
Coming back to the high temperature regime, at least around 300 K, nearestneighbor hopping is predominant. This leads to the question whether the charging
energy of such cluster materials can be chemically tailored by means of bifunctional spacer molecules, which define the inter-particle spacing by the molecule
length. This concept has been realized in the family of ligand-stabilized metal
clusters Au55 (PPh3 )12 C6 and Pd561 (phen)36 O200 by using bifunctional spaces with
NH2 - or SH-groups on each terminus [181–183]. While the cluster size is kept
constant when the cluster spacing is increased by the spacer molecules, an almost
linear increase in the activation energy was observed. This can be explained by the
decrease in the junction capacitance C, which scales with 1=D (D ¼ particle spacing) as long as the other geometric and dielectric parameters are kept constant.
This relationship is valid as long as the spacer molecules are not covalently
bound to the cluster surface and the spacer molecules have no delocalized pelectron system along their backbone between the termini. As soon as covalently
linking species equipped with delocalized p electrons enable inter-cluster electron
transfer to occur, the activation energy drops, depending on the electronic structure
of the molecule and its length. This shows that, at least by integral conductance
measurements, the electron transport through individual organic molecules can be
revealed, which is supporting information about the nature of electron transport
on the molecular level.
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6
Biomaterial-Nanoparticle Hybrid Systems:
Synthesis, Properties, and Applications
Eugenii Katz, Andrew N. Shipway, and Itamar Willner
6.1
Introduction
The unique electronic, optical and catalytic properties of metal and semiconductor
nanoparticles (1–200 nm), together with the different methods available for the
preparation of nanoparticles of controlled shape and size, provide exciting building
blocks for nanoscale assemblies, structures and devices. Nobel Laureate Richard
Feynman in his visionary lecture ‘‘There is plenty of room at the bottom’’ [1]
inspired the concepts for the rapidly exploding research topic of nanotechnology.
Although the term ‘‘nanotechnology’’ had not appeared on the horizon, Feynman
said: ‘‘What I want to talk about is the problem of manipulating and controlling
things on a small scale . . . What I have demonstrated is that there is room – that
you can decrease the size of things in a practical way . . . I will not discuss how we
are going to do it, but only what is possible in principle . . . We are not doing it
simply because we haven’t yet gotten around to it.’’ Four decades later scientists
have learnt that the manipulation of atoms, molecules and clusters on surfaces is
feasible, and that new fundamental physics governs the properties of nano-objects.
The miniaturization of structures by conventional [2] and electron-beam lithography [3] is reaching the theoretical limits of ca. 50 nm. For the further miniaturization of chemical objects, alternative approaches must be developed. Following Feynman’s vision, one may employ atoms and molecules as building units
for the ‘‘bottom-up’’ assembly and construction of architectures with nanometer
dimensions. Nanoparticles consisting of metals [4] (e.g., Au, Ag, Pt, and Cu) or
semiconductors [5–9] (e.g., PbS, Ag2 S, CdS, CdSe, and TiO2 ) seem to be attractive
units for the engineering of such structures. The unique electrical charging properties of these particles [10], as well as their optical and photophysical features [11],
such as size-controlled plasmon absorbance, photonic electron-hole pair generation, and fluorescence, allow the addressing of particles by external electronic and
photonic signals.
A variety of synthetic methodologies for the preparation of nanoparticles within
a narrow size distribution are available [4]. Often, the nanoparticles are prepared
by ‘‘wet chemistry’’ procedures, where the clustering of the metal atoms or semiNanoparticles: From Theory to Application. Edited by Günter Schmid
Copyright 8 2004 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
ISBN: 3-527-30507-6
6.1 Introduction
conductor molecules proceeds in the presence of a surface capping ligand. This
capping ligand binds to the metal/semiconductor clusters, prevents aggregation of
the particles into bulk material, and controls the final dimensions of the nanoparticles. Many capping systems are available, including hydrophobic monolayers
[12], positively or negatively charged monolayers [13], and polymer layers [14].
Association of molecular units to the nanoparticles introduces chemical functionalities that can provide recognition or affinity interactions between different
appropriately modified particles, and thereby dictate structure when aggregation
occurs [15]. New collective properties of aggregated particles [11] such as coupled
plasmon absorbance, interparticle energy transfer, and electron transfer or conductivity may be observed in the clustered assemblies.
The chemical functionalities associated with nanoparticles enable the assembly
of 2D and 3D nanoparticle architectures on surfaces [11]. Composite layered or
aggregated structures of molecule- or macromolecule-crosslinked nanoparticles
on surfaces have been prepared, and the specific sensing of substrates [16], tunable
electroluminescence [17], and enhanced photoelectrochemistry [18] have been
accomplished. The assembly of nanoparticle architectures on surfaces has also led
to the fabrication of nanoscale devices such as single electron transistors [19],
nanoparticle-based molecular switches [20], metal-insulator-nanoparticle-insulatormetal (MINIM) capacitors [21], and others. Several reviews have addressed recent
advances in the synthesis and properties of nanoparticles [11, 22–24] and the progress in the integration of composite nanoparticle systems with surfaces [11, 25–27].
The conjugation of nanoparticles with biomaterials is a tempting research project that provides a route into nanobiotechnology. Evolution has optimized fascinating macromolecular structures exhibiting unique recognition, transport, and
catalytic properties. The conjugation of nanoparticles with biomaterials could
provide electronic or optical transduction of the biological phenomena. Enzymes,
antigens and antibodies, and receptors have dimensions in the range of 2–100 nm,
comparable to those of nanoparticles, and thus the two have structural compatibility. Several fundamental features show biomaterials to be important future building blocks for nanoparticle architectures:
1. Biomaterials reveal specific and strong complementary recognition interactions,
e.g., antigen-antibody, nucleic-acid-DNA, hormone-receptor. The functionalization of a single kind of nanoparticle or different kinds of nanoparticles with
biomaterials could lead to biomaterial-nanoparticle recognition and thus to selfassembly.
2. Various biomaterials include several binding sites, e.g., the two Fab-sites of
antibodies, the four binding domains of streptavidin or concanavalin A. This
allows the multidirectional growth of nanoparticle structures.
3. Proteins may be genetically engineered and modified with specific anchoring
groups. This facilitates the aligned binding to nanoparticles, or the site-specific
linkage of the biomaterial to surfaces. Consequently, directional growth of nanoparticle structures may be dictated. Furthermore, other biomaterials, such as
double-stranded DNA, may be synthetically prepared in complex rigidified
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structures that act as templates for the assembly of nanoparticles by intercalation, electrostatic binding to phosphate groups, or association to functionalities
tethered to the DNA.
4. Enzymes provide catalytic tools for the manipulation of biomaterials. For example, the ligation of nucleic acids or the endonuclease scission processes of
nucleic acids provide effective tools for controlling the shape and structure of
biomaterial-nanoparticle hybrid systems. In this context, it is important to note
that Mother Nature has developed unique biocatalytic replication processes. The
use of biocatalysts for the replication of biomaterial-nanoparticle conjugates
may provide an effective system for the formation of nanostructures of predesigned shapes and compositions.
It is the aim of this chapter to review the synthesis of biomaterial-nanoparticle
hybrid systems, as well as the organization of these systems as functional devices
(Figure 6-1). We describe their properties, the methods of assembling twodimensional and three-dimensional nanoparticle structures on surfaces, and the
operation of these biomaterial-nanoparticle structures as functional devices. Specifically, we discuss the use of the systems as sensing devices, as building blocks
for electronic circuitry, and as electronic and optoelectronic elements. It is our aim
to highlight the fact that the scientific accomplishments of the past few years have
already addressed important aspects of nanobiotechnology.
The conceptual generation of nanoparticle-biomaterial
conjugates and their assembly to give functional devices.
Fig. 6-1
6.2 The Synthesis and Properties of Biomaterial-Functionalized Nanoparticles
6.2
The Synthesis and Properties of Biomaterial-Functionalized Nanoparticles
Biological molecules have been immobilized on polymer matrices and inorganic
supports by a variety of techniques, including physical adsorption, electrostatic
binding, specific recognition and covalent coupling [28–32]. These supports,
modified with biological molecules such as proteins/enzymes, antigens/antibodies,
and DNA/oligonucleotides, have been used for numerous biotechnological applications, including affinity separations, biosensing, bioreactors, and biofuel cells.
Recently, these immobilization techniques – developed for the functionalization
of macro-size supports – have been applied to bring together biomolecules and
nanoparticles.
6.2.1
Functionalization by Electrostatic Adsorption
The simple adsorption of biomolecules on nanoparticles has frequently been used
[28–30] and studied, these biomolecules ranging from low-molecular-weight organic substances (e.g., vitamin C [33]) to large protein/enzyme molecules [34–36].
In the case of nanoparticles that are stabilized by anionic ligands such as carboxylic
acids (citrate, tartarate, lipoic acid), the adsorption of proteins originates from
electrostatic interactions (Figure 6-2) [37]. For example, gold and silver nanoparticles produced by citrate reduction have been functionalized with immunoglobulin
(IgG) molecules at pH values slightly above the isoelectric point of the citrate ligand [38]. This allows effective binding between the positively charged amino acid
side chains of the protein and the negatively charged citrate groups of the colloids.
Other examples of protein coating through electrostatic interactions include the
direct adsorption of heme-containing redox enzymes at citrate-stabilized silver
nanoparticles [34, 39, 40] and the binding of basic leucine zipper proteins to lipoic
acid-stabilized semiconductor CdSecore /ZnSshell particles [41]. Differential scanning calorimetry (DSC) has been used to demonstrate the electrostatic nature of
proteins adsorbed on silica particles, and allows the calculation of thermodynamic
parameters [42]. The electrostatic deposition of biomolecules, particularly proteins
or enzymes, can also be extended to a multi-layer level [43]. Protein molecules
Fig. 6-2
The assembly of nanoparticle-protein conjugates by electrostatic interactions.
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
electrostatically attracted to the charged nanoparticles can provide a base interface
for the further deposition of an oppositely charged polyelectrolyte polymer that
again allows the deposition of a secondary protein layer. Multilayer films of bovine
serum albumin (BSA) [44], immunoglobulin G (IgG) [44], b-glucosidase (b-GLS)
[45], glucose oxidase (GOx) [46, 47], urease (Ur) [48], and horseradish peroxidase
(HRP) [47] have thus been assembled on polystyrene nanoparticles by the alternate
deposition of the protein and oppositely charged polymers. The protein/polymer
multilayer shell could be varied from several to hundreds of nanometers in thickness. This strategy permits the preparation of functional films with a high density
of enzyme molecules on nanoparticles.
Similarly, the electrostatically driven adsorption of negatively charged DNA molecules on positively charged CdS nanoparticles has been studied in detail [49–53].
An increase in the loading of nanoparticles of organic materials, upon layer-bylayer deposition, increases the sensitivity of analytical methods when these particles are utilized. For example, polystyrene particles were functionalized with a
multilayer assembly of polyelectrolytes functionalized with a fluorescent dye. The
assembly was terminated with an IgG layer that allowed the use of the modified
particles as labels in fluorescence immunoanalysis [54]. The increase of the loading of the particles with the dye because of multilayer deposition resulted in higher
fluorescence output, and thus in a higher sensitivity of the analysis.
6.2.2
Functionalization by Chemisorption of Thiol-Derivatized Biomaterials
In some cases, strong chemisorption of proteins on Au nanoparticles can originate
from the binding of thiol groups (from cysteine residues) existing in the proteins
(e.g., immunoglobulins (IgG) [38], serum albumin [55]) to the Au surface (Figure
6-3A). If no such residues are available in the native proteins, thiol groups can be
incorporated by chemical means (e.g., with Traut’s reagent, 2-iminothiolane (1) in
Figure 6-3B) [56] or by genetic engineering. Some proteins and enzymes preserve
The formation of nanoparticle-protein conjugates by
the adsorption of nanoparticles on A natural, and B synthetic
thiol groups of the protein.
Fig. 6-3
6.2 The Synthesis and Properties of Biomaterial-Functionalized Nanoparticles
their native structures and activities when they are directly adsorbed on nanoparticles. For instance, a surface-enhanced resonance Raman (SERR) spectrum
of native hemoglobin was obtained when it was adsorbed on citrate-stabilized Ag
colloids. Reversible dioxygen and carbon monoxide binding and the R- to T-state
transition were also observed [36]. A non-heme iron enzyme, chlorocatechol dioxygenase, physically adsorbed on citrate-reduced Ag colloid, also showed its native
structure and biocatalytic activity upon application of SERR spectroscopy [34].
However, physically adsorbed biomolecules can be readily lost from the surface,
and adsorbed proteins are often prone to denaturation, thereby losing their biocatalytic or biorecognition activities. Careful investigations of some proteins and
enzymes directly adsorbed on nanoparticles demonstrate their conformational
changes and loss of biological activity. For example, circular dichroism (CD), DSC,
and fluorescence spectroscopy applied to bovine fibrinogen adsorbed on TiO2
nanoparticles indicate that the a-helix content of the protein decreases markedly
upon their adsorption [35].
Nucleic acids can be synthesized with tethered alkanethiol groups by using
appropriate phosphoramidate precursors in a solid-phase synthesis. The nalkylthiolated nucleic acids have been used extensively in the preparation of
DNA-functionalized gold nanoparticles [57, 58] and CdSecore /ZnSshell [59, 60] or
CdS [52] semiconductor nanoparticles. DNA molecules functionalized with a
steroid-disulfide derivative can readily bind to Au nanoparticles and demonstrate
higher adsorption stability because two sulfur atoms of the disulfide anchor group
are involved in the attachment process [61]. Enhanced binding of nucleic acids to
metal nanoparticles may be accomplished by the use of oligonucleotides that
include several adenosyl phosphothiolate residues at their ends, thus enabling
multi-point attachment to the nanoparticles [62].
6.2.3
Functionalization by Specific Interactions
Nanoparticles functionalized with groups that provide sites for the affinity binding
of biomolecules have been used for the specific attachment of proteins and oligonucleotides. For example, streptavidin-functionalized Au nanoparticles have been
used for the affinity binding of biotinylated proteins (e.g., immunoglobulins and
serum albumins [63]) or biotinylated oligonucleotides (Figure 6-4A) [64, 65]. Likewise, biotinylated CdSe/CdS/SiO2 nanoparticles can be bound to streptavidin [66].
Also, nanoparticle-antibody conjugates were used for affinity binding of their respective antigens (Figure 6-4B) [67]. This may be advantageous, as antibodies associated with nanoparticles can demonstrate association constants with respective
antigens that are even higher than those of the free antibody [68].
By covalently attaching proteins to nanoparticle surfaces, problems of instability,
reversibility and inactivation can be overcome [43]. Low-molecular bifunctional
linkers that have anchor groups for attachment to nanoparticle surfaces and functional groups for further covalent coupling to the target biomolecules were extensively used for generating covalent conjugates of biomolecules and various nano-
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The assembly of nanoparticle-biomaterial conjugates by the use of bioaffinity interactions, A by the use of
streptavidin-biotin interactions, and B by the use of antibodyantigen interactions.
Fig. 6-4
particles [69]. Anchor groups such as thiols, disulfides, or phosphane ligands are
often used for the binding of the bifunctional linkers to Au, Ag, CdS, and CdSe
nanoparticles. These anchor groups readily substitute weakly adsorbed molecules
stabilizing the nanoparticles, or may be incorporated in the nanoparticle synthesis,
resulting in a nanoparticle surface providing functional groups for further reactions. Alkoxysilane or halogen-silane groups are used for the covalent attachment
of bifunctional linkers to the surfaces of SiO2 and other oxide-coated nanoparticles.
A wide variety of terminal functional groups is available in different bifunctional
linkers: amino, active ester, and maleimide groups are used for the covalent coupling of biological compounds by means of carbodiimide-mediated esterification
and amidation, or reaction with thiol groups [69]. Single functional groups (e.g.,
active ester or amino units) on Au nanoparticles are available either by a statistical
approach or by the modification of chemically defined clusters. These structures
provide unique synthetic routes for the covalent binding of a single target biomolecule per nanoparticle [70–72].
6.2.4
Properties of Nanoparticle-Biomaterial Composites
The functionalization of nanoparticles with biomolecules results in changes in the
properties of the nanoparticles and their interactions with the environment. Upon
adsorption of vitamin C on TiO2 nanoparticles, the optical properties of the particles were red shifted by 1.6 eV. This is the result of charge transfer, and originates
from the specific binding of the electron-donating modifier to corner defects on the
surface of the nanoparticles [33]. The solubility of nanoparticles in water can be
6.2 The Synthesis and Properties of Biomaterial-Functionalized Nanoparticles
The electrochemically-controlled recognition of flavin
by a diamidopyridine-functionalized nanoparticle.
Fig. 6-5
greatly improved by the functionalization of their surfaces with highly hydrophilic
biomolecules [73]. While Au nanoparticles modified with long chain alkanethiols
are only soluble in low-polarity organic solvents, Au nanoparticles with o-carboxylic acid-functionalized alkanethiolate ligand shells are soluble in polar organic
solvents such as ethanol and acetone. Au nanoparticles modified with biomolecules such as thiopronin or coenzyme A demonstrate excellent solubility in water.
Alteration of the chemical properties of the biomolecules covering nanoparticles
by external signals (e.g., electrical, optical) can be used to control interactions of
the modified nanoparticles with the environment, for example, to control the
binding of a secondary modifier or the aggregation of the nanoparticles. Complex
formation between a diaminopyridine derivative and a flavin derivative was studied
at an Au nanoparticle surface [74]. It has been shown that more stable hydrogen
bonding is provided when the flavin derivative is electrochemically reduced (Figure
6-5). This allows electrochemically controlled switching of the binding of the bioorganic molecules to the organic-functionalized shell of the nanoparticle.
The interactions of biomaterial-functionalized nanoparticles with other functionalized nanoparticles or other biomaterials such as proteins or DNA molecules
provided by affinity properties of the biomaterial shell components are described in
the following sections.
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6.3
Biomaterial-Functionalized Nanoparticles for Controlled Chemical Reactivity
The interactions of functionalized nanoparticles with biomaterials or within biomaterial structures can control the chemical reactivity of the biomolecules. Alternatively, they can report on the state of the reaction or reactants, allowing control to
be exercised externally. It is well known that small molecules and polymers can
affect the chemical reactivity of biomolecules. If there are several possible parallel
reactions, the effect produced by a promoter/inhibitor on a specific chemical reaction can change the effective chemical path of the whole process, resulting in a
regulation of the biochemical system. Molecular labels (e.g., fluorescent dyes)
incorporated into biomolecules can report on the state of the biomolecule, transducing the molecular operation into an output signal used, for example, in
biomaterial-based diagnostics. Functionalized nanoparticles can operate in the
same way, demonstrating properties of a biopromoter/bioinhibitor or a reporter,
with some advantages over their molecular-sized counterparts due to the unique
photophysical and electronic properties of the nanoparticles.
6.3.1
Controlling DNA Reactivity
Biocatalytic reactions occurring at DNA molecules, such as transcription or translation, require the association of enzymes with the DNA chain. The competitive
association of other molecules to the DNA chain can decrease the enzyme binding
and thus inhibit the process, finally resulting in gene regulation. It has been
shown that small molecules [75], dendrimers [76], and polymers [77] are capable
of efficient binding to DNA that in some cases leads to inhibition of the transcription machinery. In these systems, however, the creation of suitably preorganized
scaffolds for the controlled binding to DNA presents a significant design challenge.
For small molecules, this challenge arises from the detailed synthesis required for
a rigid scaffold containing several functional groups. For polymer systems, the
precise placement of substituents along the backbone as well as the control of
the polymer macro-conformation in solution are difficult to achieve. An excellent
solution to these problems is the application of ‘‘mixed monolayer-protected Au
clusters’’ (MMPCs) [78]. The self-assembled monolayer coating of the Au nanoparticles presents a highly organized surface for the recognition of DNA that is of a
similar size scale (6–10 nm) to that of DNA-binding proteins. The central metal
core rigidifies the particle, limiting the organic components to a much smaller
subset of structures than a similarly sized polymer. The mobility of the thiols on
the nanoparticle surface also allows the formation of recognition elements on the
nanoparticle surface to be controlled in subsequent steps [79]. Nanoparticles for
the recognition of double-stranded (ds) DNA were prepared by the partial displacement of octanethiol on the Au surface with 11-mercaptoundecyltrimethylammonium [80]. The mixed monolayer of this MMPC system is designed to interact
with DNA based on the electrostatic attraction of the positively charged quater-
6.4 The Aggregation of Biomaterial-Functionalized Nanoparticles
nary ammonium salt and the negatively charged DNA phosphate groups. These
MMPCs bind to the ds-DNA in competition with the enzymes involved in gene
regulation, thus resulting in effective inhibition of the biocatalytic transcription
process.
6.3.2
Reporting on DNA Reactivity
DNA molecules functionalized at specific positions with a fluorescent dye and a
quencher can produce a fluorescence output controlled by the distance separating
the dye and the quencher, thus reporting on the conformation of the DNA chain.
Au nanoparticles can quench fluorescence of chromophores to an extent that is
100-fold better than molecular quenchers, and they also have higher quenching
efficiency for dyes emitting in the near infrared region. A DNA oligonucleotide was
functionalized with a fluorescent dye at the 3 0 end, and an Au nanoparticle functionalized with a single maleimide group was covalently linked to the 5 0 end
through a (CH2 )6 -SH group [81]. The molecule could adopt two conformations
(Figure 6-6): a stem-loop structure with the fluorophore and the Au nanoparticle
held in close proximity (hairpin state), and a rod-like structure with them far apart
(open state). The hairpin state (2) is self-generated because of the intramolecular
complementarity of the terminal parts of the DNA chain that make a ds-structure
within the same molecule, whereas the open state (3) is produced upon interaction
and hybridization with an analyte DNA molecule (4) that exhibits complementary
to the central part of the DNA chain. This target DNA hybridizes with the modified
DNA, resulting in the formation of ds-DNA, thus opening the hairpin, increasing
the distance between the fluorophore and the nanoparticle, and thereby enhancing
the fluorescence. The process and the reporting signal only occur in the case of
complementarity between the analyte DNA and the central part of the sensing
DNA. The system was thus successfully applied for single-mismatch detection in
DNA sequences [81].
6.4
The Aggregation of Biomaterial-Functionalized Nanoparticles
The organization and patterning of inorganic nanoparticles into two- and threedimensional (2D and 3D) functional structures is a fundamental prerequisite for
the assembly of chemical, optical, magnetic and electronic devices [11]. Many approaches have been described for the formation of 2D and 3D arrays of metal and
semiconductor nanoparticles, for example, using (i) solvent evaporation of hydrophobic colloids [82–84], (ii) random inclusion of the nanoparticles into gels and
glassy matrices [85], (iii) template-directed synthesis at structured surfaces [86]
in porous protein crystals [87, 88] or bacterial superstructures [89], and (iv) chemical coupling in solution by means of bivalent crosslinker molecules [90–92]. An
example of the latter is the alkanedithiol-directed aggregation of Au nanoparticles
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
The change between hairpin (2) and extended (3)
conformations of a DNA chain by the use of a second strand
(4). In the extended conformation, the ends (which bear an Au
nanoparticle and a dye) are far apart, and in the hairpin
conformation they are close.
Fig. 6-6
[92]. The aggregation of nanoparticles induced by specific biological interactions
attracts interest as a self-assembly process for the building of complex nanostructures exhibiting new collective properties.
There are several advantages in utilizing biomaterials as building blocks of
nanoparticle structures:
6.4 The Aggregation of Biomaterial-Functionalized Nanoparticles
1. The diversity of biomaterials facilitates the selection of building units of predesigned size, shape and functionality.
2. The availability of chemical and biological means to modify and synthesize biomaterials, e.g., synthesis of nucleic acids of predesigned composition and shape,
eliciting monoclonal antibodies, or modifying proteins by genetic engineering,
pave the way to the construction of biomaterials for the assembly of nanoparticles.
3. Enzymes may act as biocatalytic tools for the manipulation of the biomaterials.
4. Hydrolysis of proteins, scission or ligation of DNA or replication of nucleic
acids may be employed as assembler tools of nanoparticle architecture through
the manipulation of the biomaterial.
5. Mother Nature has developed routes for the repair of biomaterials. These
processes may be applied for the stabilization of the biomaterial/nanoparticle
structures.
6. Crosslinking nanoparticles with enzyme units may generate biocatalytic assemblies of pre-designed functionality.
These different features of the biomaterial crosslinking units provide the flexibility for the generation of nanoparticle structures of tunable physical, chemical
and functional properties. For the generation of biomaterial-crosslinked nanoparticles, two kinds of biomaterial-functionalized nanoparticles with complementary units should participate in the assembly process. Biomaterials utilized in the
fabrication of such biomaterial-nanoparticle aggregates include biological protein
host-guest pairs such as biotin-avidin, hapten-antibody [93] and complementary
oligonucleotides [94, 95].
6.4.1
Receptor-Induced Aggregation of Guest-Functionalized Nanoparticles
Protein-based recognition systems can be used to organize inorganic nanoparticles into network-aggregates, for instance with the interaction between D-biotin
and the biotin-binding protein streptavidin (Sav) [93]. The recognition between
water-soluble biotin and the homotetrameric protein Sav is characterized by an
extraordinary affinity constant of K a > 10 14 M1 , which makes it the strongest
ligand-receptor binding interaction presently known [96]. Another great advantage of Sav is its extreme chemical and thermal stability. The applicability of the
biotin-Sav system for generating supramolecular aggregates is enhanced by the
availability of various biotin analogs [97, 98] and recombinant Sav mutants [99,
100]. A wide range of rate and equilibrium constants (K a from 10 0 to 10 15 M1 )
are available, allowing the construction of carefully designed nanoparticle aggregates. Many biotinylated materials are commercially available or can be prepared under mild conditions. Two routes for nanoparticle aggregation by biotinSav interactions are possible: (i) nanoparticles functionalized with biotin groups
can be crosslinked with the tetra-valent Sav receptor (Figure 6-7A), and (ii) nanoparticles functionalized with Sav can be crosslinked with a dibiotin derivative
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
The use of biotin-streptavidin
interactions to build nanoparticle networks, A
using streptavidin to link biotin-functionalized
nanoparticles, and B using a biotin dimer to
Fig. 6-7
link streptavidin-functionalized nanoparticles.
C The use of streptavidin-linked thioated biotin
to build nanoparticle networks.
(5) (Figure 6-7B). Several synthetic routes can by applied to functionalize metal
and semiconductive nanoparticles with biotin derivatives. In the simplest method,
thiol or disulfide biotin derivatives are directly adsorbed on metal (e.g., Au, Ag)
nanoparticles. Alternatively, nanoparticles can first be coated with an organic
‘‘shell’’ (e.g., by the polymerization of a trialkoxysilyl derivative or by polymer adsorption) and then covalently modified with biotin (e.g., by carbodiimide coupling).
Au nanoparticles functionalized with adsorbed disulfide-derivatized biotin units
(6) and then crosslinked with the added Sav have been shown to yield an aggregate
of Au nanoparticles with the biotin-Sav recognition pairs between the nanoparticles
(Figure 6-7A) [93, 101, 102]. A similar architecture can be built by reversing the
steps: Sav was interacted with the disulfide biotin derivative to produce a complex,
which was then reacted with Au nanoparticles (Figure 6-7C) [101]. In both cases,
fast, spontaneous aggregation of the Au nanoparticles was observed, resulting in a
non-ordered network of particles. Dynamic light scattering, small-angle X-ray scattering and transmission electron microscopy were used to follow the aggregation
process and to characterize the product. The kinetics of this process was fitted to
the theoretical Smoluchowski model of aggregation, modified for nanoparticles
[102]. This kinetic analysis allows control of the aggregate structure by optimiza-
6.4 The Aggregation of Biomaterial-Functionalized Nanoparticles
Fig. 6-8
Biotin-functionalized ferritin.
tion of the conditions (molar ratio of the biotin-functionalized nanoparticles and
the Sav-crosslinker). Another example of biotin-Sav-crosslinked nanoparticles involves the native protein ferritin, which consists of a hollow polypeptide shell, 8
and 12 nm in internal and external diameters, respectively, and a 5 nm diameter
ferric oxide core (5Fe 2 O3 9H2 O) [103]. The ferric oxide nanoparticle can easily
be removed from the protein shell by reductive dissolution, and the empty cage of
apoferritin can be remineralized in vitro with a range of inorganic oxides, sulfides
or selenides (e.g., MnO, FeS, CdS, CdSe) [104–107]. A succinimidyl active ester derivative of biotin (7) was covalently linked to the amino groups of 60–70 lysine residues provided by the polypeptide shell of ferritin, yielding a biotin-functionalized
protein shell containing the ferric oxide core (Figure 6-8). Addition of Sav resulted
in the crosslinking and therefore the aggregation of the functionalized ferric oxide
nanoparticles [108]. Considering the library of core nanoparticles possible, this
ferritin-based aggregate is a very promising material for producing electronic,
magneto-electronic, and optoelectronic devices.
Nanoparticle aggregates have also been produced by means of the interaction
between antigens and antibodies [93, 109]. In the simplest approach, anti-DNP
IgE (DNP, dinitrophenyl) antibodies were adsorbed on 12 nm Au nanoparticles
and were then crosslinked with a synthetic antigen (bis-N-2,4-dinitrophenyloctamethylene diamine) comprising two hapten DNP head groups separated by a
spacer. A more sophisticated approach allows the association of two different kinds
of nanoparticles. For this purpose, two different nanoparticles functionalized with
different kinds of bioreceptors (e.g., different antibodies or an antibody and Sav)
should be crosslinked with a ligand containing two different head groups providing affinity binding to the two respective receptors. For example, two kinds of
Au nanoparticles were functionalized, one with anti-DNP IgE and another with
Sav [93, 109]. These were crosslinked with a linker (8) containing a biotin unit and
a DNP group separated with a spacer (Figure 6-9). Each ligand recognizes its
respective receptor (biotin binds to Sav, DNP binds to anti-DNP IgE), thus creating
a network composed of both nanoparticles.
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
Formation of a network composed of two kinds of
nanoparticles functionalized with different recognition units
and crosslinked with a bifunctional linker.
Fig. 6-9
A novel approach to nanoparticle crosslinking involves the application of specially designed de novo proteins with specific affinities to inorganic crystal surfaces
[69]. These structures can specifically bind directly to unmodified surfaces of various semiconductive nanoparticles, thus resulting in their bridging and aggregation. Based on a combinatorial library of about 10 9 random 12-mer peptides, de
novo proteins with specific binding capabilities to single-crystal semiconductors
including GaAs(100), GaAs(111), InP(100), and Si(100) have been discovered [110].
6.4.2
Nucleic Acid-Functionalized Nanoparticles for Controlled Aggregation
DNA molecules are very powerful and versatile linkers for the controlled aggregation of nanoparticles [111, 112]. DNA chains are of a controllable length and composition, providing predictable properties and enormous control over complementarity. ds-DNA provides rigid spacers that allow control over the distance between
nanoparticles. Long DNA chains can be utilized as templates ready to accommodate complementary-DNA-functionalized nanoparticles in a controlled order and
with controlled distances between the nanoparticles.
In order to demonstrate the DNA-mediated binding of nanoparticles, two batches
of 13 nm diameter Au nanoparticles were separately functionalized with two thiolderivatized non-complementary DNA oligonucleotides: 3 0 -thiol-TACCGTTG-5 0 (9)
adsorbed on nanoparticle A and 5 0 -AGTCGTTT-3 0 -thiol (10) adsorbed on nanoparticle B (Figure 6-10A, see p. 384) [113]. A DNA linker (11) was constructed,
composed of ds-DNA in the center and two single-stranded fragments on the ends.
These single-stranded sections were complementary to the oligonucleotides connected to the nanoparticles. Addition of the linker to a mixed solution of the two
6.4 The Aggregation of Biomaterial-Functionalized Nanoparticles
DNA-functionalized Au nanoparticles results in crosslinking of the nanoparticles.
The resulting nanoparticle aggregate should theoretically have some degree of
regularity, but this could not be observed, as both nanoparticles had identical cores.
Aggregation changes the color of the solution because of a collective phenomenon
originating from the short-distance interparticle separation. Separate Au nanoparticles have a red color as a consequence of their plasmon absorbance, but when
they are aggregated additional absorbances arise at lower energies because of
interparticle plasmon coupling (Figure 6-10B). This color change can be used as a
simple test for the recognition between complementary DNA sequences associated
with the linker and the Au nanoparticles. The association process is temperaturedependent, and the aggregated Au nanoparticles can dissociate upon elevation of
temperature and associate again upon decreasing the temperature, thus resulting
in the reversible changes of the spectrum (Figure 6-10C). In a control experiment,
a non-complementary linker was added to the mixed particle solution, and no color
changes were observed.
In a more sophisticated investigation, a mixture of DNA-functionalized Au
nanoparticles of diameters 40 nm (A) and 5 nm (B) was crosslinked with a DNA
linker [114]. The larger nanoparticle (A) accommodates many small nanoparticles
(B) around it, producing a so-called ‘‘satellite’’ ABn system (Figure 6-10D). Interconnections between these ‘‘satellite’’ aggregates yield a 3D-network (ABn )m . An
important factor controlling the network optical properties is the length of the
DNA oligonucleotide: it controls not only the interparticle distances, but also the
rate of aggregation and thus the number of nanoparticle per aggregate [115]. In
a similar approach, DNA-functionalized CdSe/ZnS nanoparticles and Au nanoparticles were crosslinked with a complementary linker to yield a binary network
[59]. Fluorescence and absorption spectroscopy indicated cooperative optical and
electronic phenomena within the resulting material. Changes in the optical and
electronic properties of nanoparticles upon their aggregation provide a way to
detect the association process, thus allowing DNA-biosensor construction [116–
119]. The application of DNA-functionalized Au nanorods instead of spherical
nanoparticles results in microsized rods composed of many crosslinked nano-units
(Figure 6-11, see p. 385) [120]. This material offers unusual anisotropic properties.
Long DNA molecules capable of accommodating several complementary DNA
fragments in different domains on the strand can serve as templates for the
construction of specific linear assemblies [121, 122]. In order to prevent the
microscopic crosslinking of nanoparticles resulting in a network, Au nanoparticles functionalized with a single maleimide group (12) were reacted with thiolderivatized DNA oligonucleotides (13). The resulting mono-DNA-functionalized
Au nanoparticles (14) were associated with the respective complementary domains
on the DNA template molecule (Figure 6-12). Individual placement of the nanoparticles at specific complementary domains of the template allows control of
distances between them, as well as controlled placement of different kinds of
nanoparticles (different sizes, different materials) on the template, which results
in a variety of properties of the nanoparticles assemblies. For example, two DNAfunctionalized Au nanoparticles can be placed on the DNA template with three
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
Fig. 6-10 A The aggregation of oligonuleotidefunctionalized nanoparticles by the addition of
a DNA strand with sections complementary to
the DNA on each nanoparticle. B Cuvettes with
the Au nanoparticles and the added DNA
strand responsible for the assembly process.
Left cuvette, at 80 C with DNA-functionalized
nanoparticles in the unhybridized state; center,
after cooling to room temperature resulting in
the nanoparticle association but before the
precipitate settles; and right, after the
associated nanoparticles settle to the bottom
of the cuvette. C Variations in the absorption
spectrum of DNA-linked nanoparticle networks
as a function of temperature. The changes
reflect different degrees of aggregation, which
is controlled by DNA association (low temperature) and dissociation (high temperature).
D A ‘‘satellite’’ system consisting of two
different sizes of nanoparticles. (Parts A–C are
adopted from [113], Figures 1–3; Part D is
adopted from [114], Figure 2, with permission.)
different orientations: ‘‘head-to-head’’ (15), ‘‘head-to-tail’’ (16), and ‘‘tail-to-tail’’ (17),
thus providing different distances between the nanoparticles.
6.4.3
Composite Assemblies of Nucleic Acids, Proteins, and Nanoparticles
A combination of the synthetic methods to prepare biomaterial-functionalized
nanoparticles (cf. Sections 6.4.1 and 6.4.2) and the application of proteins with
affinity properties (e.g., Sav, antibodies), oligonucleotides and nanoparticles as
building blocks, allows the construction of very sophisticated biomolecular/
nanoparticle clusters (Figure 6-13A) [69]. In order to functionalize Au nanoparticles with antibodies, a multi-step procedure was performed [123]. 5 0 -Thiolmodified oligonucleotide (18) was covalently bound to Sav using the heterobifunctional crosslinker sulfosuccinimidyl-4-(maleimidomethyl)cyclohexane-1-carboxylate
6.4 The Aggregation of Biomaterial-Functionalized Nanoparticles
Fig. 6-11 Transmission electron micrographs of gold nanorods
organized by DNA hybridization ((a) and (b) different
magnifications). (Reproduced from [120], Figure 1, with
permission.)
Fig. 6-12
The assembly of controlled multiparticle composites by the use of DNA hybridization.
(19). The resulting conjugate has four binding sites for biotin ligands and one nucleic acid function. The Sav receptor sites were reacted with biotinylated antibodies
(20) directed against mouse IgG (aM-IgG) or rabbit IgG (aR-IgG), resulting in
DNA-tagged antibody conjugates that can bind to the complementary DNA from
one side and to the respective antigens from the other. This architecture was reacted with Au nanoparticles functionalized with oligonucleotides complementary
to the DNA tag (21), thus finally resulting in the Au nanoparticles modified with
the antibody molecules. The resulting conjugate was then applied as a label for the
immunoassay of the respective antigens. The linkage of the Au nanoparticle to the
antigen-antibody complex was further amplified by the electroless deposition of Ag
on the Au nanoparticles (Figure 6-13B) [123].
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
Fig. 6-13 A The formation of an Au nanoparticle-IgG
conjugate, and B its use in the sensing of antigens.
A combination of receptor proteins and oligonucleotides as binding blocks for
the directed aggregation of nanoparticles allows the complex structuring of the
assemblies. For example, Sav was interacted with biotinylated DNA, resulting in a
four-DNA-Sav conjugate [124]. Application of this conjugate as a building block
in the aggregation of DNA-functionalized Au nanoparticles yields a network with a
motif different from that resulting from direct aggregation of DNA-functionalized
nanoparticles bound by the complementary oligonucleotide. Even more fascinating aggregation was achieved when Au nanoparticles were primarily associated
with Sav to yield nanoclusters that were then placed on a DNA template [125]
(Figure 6-14). In this work, 1.4 nm Au nanoparticles with a single amino substituent were derivatized with a biotin group, and the biotin moiety was used to
organize the nanoclusters into the tetrahedral superstructure defined by the ge-
6.4 The Aggregation of Biomaterial-Functionalized Nanoparticles
Fig. 6-14 The construction of four-nanoparticle clusters and
their assembly into linear composites along a DNA strand.
ometry of the biotin-binding sites of the Sav. Prior to cluster formation, Sav was
derivatized with a single DNA molecule. The resulting tetrahedral four-Au nanoparticle clusters (22), each with a single DNA tail, were assembled on a DNA template molecule (23) with positions and distances dictated by the sequence of the
DNA template. This approach could allow the formation of almost any imaginable
structure composed of various nanoparticles of different sizes and materials.
An interesting application of submicron metallic ‘‘barcodes’’ for the analysis of
biomaterials such as DNA or antigens has recently been reported [126]. ‘‘Barcodes’’
a few hundred nanometers across and several microns long composed of many
different metals (Pt, Pd, Ni, Co, Ag, Cu and Au) were prepared by the electrochemical deposition of the metals in a porous Al2 O3 membrane (Figure 6-15A).
The structure of the barcode could be identified by optical reflectivity or field
emission-SEM. Figure 6-15B shows an optical image of an AuaAg multistripebarcode consisting of 240, 170, 110 and 60 nm Ag stripes separated by ca. 550 nm
Au stripes. It was suggested that with the available optical resolution, the number
of possible readable barcodes consisting of two metals with a coding length of 6.5
mm is 4160, and for three-metal barcodes the 8:0 10 5 distinctive striping patterns
are possible. The use of a mixture of ‘‘barcodes’’, each used to identify a specific
biomaterial adsorbed on it, could enable the simultaneous analysis of numerous
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
Fig. 6-15 A The construction of mixed metal ‘‘barcodes’’ in an
alumina membrane. B An optical image of an Au-Ag multistripe
‘‘barcode’’ nanorod. (Adopted from [126], Figures 1 and 3, with
permission.)
components in biological mixtures. Thus, the barcode concept adds a nanoparticlebased analysis method that competes with chip-arrays and encoded microbeads. In
addition, as the tag identification can be performed optically, fluorescence can be
used to characterize a biorecognition event. For example, Figure 6-16A outlines
the fluorescence detection of an analyte DNA (24) by a primer nucleic acid
(25) linked to the Ag stripes of an AgaAuaAg barcode. A fluorophore labeled
(TAMARA dye) nucleic acid (26) hybridizes to the free segment of the analyte
DNA. Figure 6-16B shows the fluorescence image (i) and reflectivity image (ii) of
the fluorescence-labeled DNA complex confined to the AgaAuaAg barcode. The
former characterizes the binding event, and the latter identifies the ‘‘barcode’’ and
thus the attached DNA probe. When the analyte DNA (24) is omitted, a very low
fluorescence background is observed (iii), and the particles are visible only in the
reflectivity image (iv). The simultaneous analysis of different biological analytes is
possible with the immunoassay of human immunoglobulin G (IgG) and rabbit
IgG. Two different barcodes, an AuaAgaAu (length 4 mm) and an AuaNiaAu
(length 8 mm) were employed, where the Ag stripe of the AuaAgaAu barcode
was modified with the anti-human-IgG antibody (27), and the Au stripes of the
AuaNiaAu barcode are derivatized with the anti-rabbit-IgG antibody (28) (Figure
6-16C). Challenging of the mixture of barcodes with the mixture of antigens followed by the labeling of the antibodies with fluorescein-labeled anti-human-IgG
6.4 The Aggregation of Biomaterial-Functionalized Nanoparticles
Fig. 6-16 A Method of using the ‘‘barcodes’’
to detect DNA. B Fluorescence (i) and
reflectivity (ii) images of ‘‘barcodes’’ with
bound fluorescent DNA; fluorescence (iii) and
reflectivity (iv) images of ‘‘barcodes’’ with
omitted analyte DNA. C Method of using the
‘‘barcodes’’ for multiple immunoassay. D
Reflectivity image (i), fluorescent image
produced by fluorescein-labeled anti-humanIgG (ii), and fluorescence image produced by
Texas Red-labeled anti-rabbit-IgG (iii).
(Adopted from [126], Figure 5, with
permission.)
(29) and Texas Red-labeled anti-rabbit-IgG (30) allows the optical detection of the
respective barcodes and the fluorescence detection of the two antigens. Figure
6-16D (i) shows the image of the two kinds of barcodes by optical reflectance, while
Figure 6-16D (ii) shows the specific fluorescence from the AuaAgaAu barcodes,
whereas Figure 6-16D (iii) shows the specific Texas Red fluorescence from the
AuaNiaAu stripes. Thus, ‘‘solution arrays’’ of highly encodable and chemically
diverse nanoparticles provide an interesting approach for the future analysis of
complex biological samples (e.g., genomics, proteomics).
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6.5
Assembly of Biomaterial-Nanoparticle Architectures on Surfaces
The immobilization of biomaterials on transducers such as electrodes, piezoelectric crystals, and field-effect transistors attracts research efforts directed to the
development of biosensors and bioelectronic devices [127–130]. Several methods
for the surface functionalization of electronic elements have been developed, and
means to couple the active sites or recognition processes of biomaterials with the
electronic elements, giving electronic transduction of the biological events, have
been demonstrated. Biosensors [131–133], biofuel cell elements [134–139] and optobioelectronic systems [140, 141] have been constructed. Similarly, the deposition
of biomaterials on non-conductive surfaces (e.g., plastics, glass), and especially the
patterning of such surfaces with biomaterials [130, 142, 143], is of high interest in
tailoring protein-chips and DNA-chips for proteomic or genomic analyses. Also,
diverse methods for the assembly of 2D and 3D nanoparticle structures on surfaces
have been developed. Various applications of these systems in the fabrication of
photoelectrochemical cells (e.g., solar cells), light-emitting diodes, electrochromic
systems, and computing devices, and also the organization of different nanometric
single-electron devices, have been discussed [11]. The integration of nanoparticlebiomaterial hybrid systems with surfaces paves the way for the generation of ordered architectures with new functionalities. The unique photonic properties of
nanoparticles may be employed to detect and probe biological recognition events
on surfaces. The spectroscopic properties of molecular adsorbates on nanoparticle
surfaces such as surface plasmon resonance (SPR), surface-enhanced Raman
spectroscopy (SERS), and surface-enhanced fluorescence may be used to probe
biomaterial structures on surfaces. Also, the unique electronic (e.g., conductive)
and catalytic properties of nanoparticles may be used to provide electron transfer
from or to redox-centers of enzymes, or to enhance chemical reactivities of or on
the biomaterial structures.
From another direction, biomaterial constituents on surfaces provide unique
features for the formation of the hybrid biomaterial-nanoparticle structures.
Biomaterials (e.g., nucleic acids, streptavidin-biotin antigen-antibody complexes)
assembled on surfaces can provide templates for the immobilization of nanoparticles. Enzymes may be used as catalytic tools for the manipulation of biomaterials on surfaces. For example, the ligation, scission or polymerization of
nucleic acids may be used to tailor nucleic acids of specific lengths and compositions. Thus, by the genetic manipulation of biomaterials, the synthesis of nucleic
acids, and the biocatalytic manipulation of the biomaterials on surfaces, complex
biomaterial-based structures on surfaces may be engineered. The incorporation
and integration of nanoparticles in biological templates can then yield new functional materials. For example, the catalytic deposition of metals on nanoparticle
seeds can generate circuitry with the shape of the biomaterial. Thus, hybrid biomaterial-nanoparticle composites on surfaces provide functional interfaces that
could be important materials for bioelectronic, electronic, optobioelectronic and
photonic applications. The following sections address the methods of organizing
6.5 Assembly of Biomaterial-Nanoparticle Architectures on Surfaces
these composite systems on surfaces and discuss recent activities that utilize these
systems for sensoric, photoelectrochemical and electronic circuitry applications.
6.5.1
Assembly of Layered Nanoparticle-Protein Assemblies
Multilayer systems assembled on solid supports and composed of cytochrome c
(Cyt c) molecules and TiO2 -nanoparticles layers have been fabricated using the
layer-by-layer deposition of the biological and inorganic components [144]. The
TiO2 -nanoparticle layers were deposited from aqueous solution by boric acid-promoted hydrolysis of hexafluorotitanate ion (liquid phase deposition process, LPD
[145]) (Figure 6-17A). The LPD technique enables the preparation of anatase TiO2
without employing organic solvents. This is essential for co-deposition of the biomaterial. The TiO2 nanoparticles produced by this technique on the surface are
negatively charged. The subsequent deposition of the positively charged protein
Fig. 6-17 A Method for the construction of TiO2 /Cyt c
multilayers on a QCM electrode. B Optical absorbance at
409 nm of assemblies of different thicknesses, showing the
buildup of Cyt c.
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
cytochrome c (isoelectric point, pI ¼ 10:1) results in electrostatic attraction between the TiO2 nanoparticles and the protein, allowing multilayer formation. The
deposition of the multilayers was followed by quartz crystal microgravimetry
(QCM), demonstrating a linear increase in the mass upon deposition of each
double layer TiO2/Cyt c, which was verified by absorbance spectroscopy (Figure
6-17B). This method could be widely extended to the preparation of other protein/
metal oxide multilayers that would facilitate the creation of photo-functional
protein/inorganic superlattices.
6.5.2
Nucleic Acid – Nanoparticle Architectures on Surfaces
As discussed in Chapter 6.4.2, nucleic acids can serve as templates capable of
binding DNA-functionalized nanoparticles at complementary segments. When
DNA templates are fixed at a surface of a solid support, the resulting assemblies of
nanoparticles can yield a pattern dependent either on the shape produced by the
DNA template itself or on the pattern produced upon its immobilization.
A natural single-stranded long-chain virus-DNA (M13mp18, 7246 bases) (31)
was hybridized with a biotinylated short-chain target oligonucleotide (32) complementary to a specific segment of the viral DNA (Figure 6-18A) [146]. It should be
noted that the template viral DNA (31) can have more than one complementary
domain, and thus is capable of accommodating more than one biotinylated-DNA
chain. The hybridized DNA-holding biotin units (33) were then reacted with
Sav-coated Au nanoparticles (34) (5 nm). The resulting negatively charged Au
nanoparticle-labeled DNA template (35) was adsorbed on a surface of positively
charged Mg 2þ -coated mica substrate and was subjected to atomic force microscopy
(AFM) (Figure 6-18B). The images obtained clearly show Au nanoparticles associated with the DNA template, yet the number of Au nanoparticles per DNA
chain was not perfectly defined, reflecting either incomplete hybridization of the
template with the biotinylated target DNA or incomplete association of the Savfunctionalized Au with the biotinylated domains. As different biotinylated nucleic
acid segments may be hybridized with the M13mp18 DNA, the entire viral DNA
may be decorated with Sav-Au nanoparticles. Indeed, such an experiment revealed
that the Au nanoparticle-functionalized viral DNA was formed on the solid support. The electrostatic association of CdS nanoparticles with DNA was applied as
the driving force for the formation of the semiconductor nanoparticle-DNA wire on
the surface [147]. Negatively charged ds-DNA molecules were located at the waterair interface containing cationic surfactant molecules (Figure 6-19). The layer at the
interface was compressed using the Langmuir-Blodgett technique in order to provide a high-density packing of the DNA molecules. Negatively charged CdS nanoparticles (3 nm) capped with thiocholine were then added to the aqueous phase,
resulting in their aggregation with the DNA. The obtained DNA/CdS aggregates
were transferred to a solid support, and electron microscopy images revealed the
formation of densely packed CdS nanoparticles along the DNA template. These
generated chains are composed of semiconductive nanoparticles with a diameter of
6.5 Assembly of Biomaterial-Nanoparticle Architectures on Surfaces
Fig. 6-18 A The controlled association of Savcoated Au nanoparticles with ss-DNA template
by means of a biotinylated ss-DNA molecule
hybridized with the template DNA. B AFM
image of a DNA strand with Au nanoparticles
specifically bound to the template through the
biotin-Sav interaction. (Part B reproduced from
[146], Figure 2(a), with permission.)
ca. 3 nm, the average distance between the centers of the adjacent nanoparticles
being ca. 3.5 nm.
The circular plasmid ds-DNA pUCLeu4 (3455 base pairs) was reacted with Cd 2þ
ions, and the resulting Cd 2þ /DNA complex was adsorbed onto a polylysine-coated
glass slide by spin coating. The complexed Cd 2þ ions were reacted with S 2 to
yield CdS, and electron microscopy images revealed the formation of CdS nanoparticles that followed the circular shape of the DNA template [148]. In another
example, positively charged lysine-capped Au nanoparticles were deposited onto a
surface coated with a negatively charged thick DNA film [149, 150]. Electrostatic
interaction between the positive charges on the Au nanoparticles and the negative
charges on the phosphate groups of the DNA template molecules leads to the assembly of nanoparticles in linear superstructures. The DNA molecules, which are
locked into a fairly rigid structure prior to addition of the lysine-capped Au nanoparticles, are rendered mobile by solvation during the addition of the nanoparticle
solution to the DNA film surface. This process facilitates reorganization into
highly regular linear nanoparticle superstructures during electrostatic complex-
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
Method of electrostatically depositing CdS nanoparticles on a DNA chain at the air-water interface using
Langmuir-Blodgett technique.
Fig. 6-19
ation. AFM and STM visualization of surface-confined DNA molecules can be
greatly improved by the addition of complementary-DNA-functionalized Au nanoparticles that not only introduce clearly visible Au clusters associated with the
DNA lying on the surface but can also change the orientation of the long DNA
molecules at the surface, producing tethered rather than flat-lying DNA chains
[151, 152]. Both these effects can facilitate the visualization of DNA molecules at
surfaces.
Negatively charged DNA molecules can complex positively charged metal ions
that are later reduced to the metallic state, producing a metal wire across the DNA
template. For example, this was accomplished by loading the DNA template with
Pd 2þ cations and adding reductive chemicals afterwards, leading to the formation
of nanoscale Pd clusters on the DNA [153]. Depending on the duration of the
reduction process, well-separated Pd clusters or a quasi-continuous Pd wire with a
grain-like structure could be produced.
A system utilizing DNA that crosslinks into a continuous network has been
investigated in an attempt to improve structural control over a longer distance
(Figure 6-20A) [154]. Three types of oligonucleotides were used, labeled 36, 37 and
38. Oligonucleotide 36 had a sequence of 17 bases, and its 3 0 -end was modified
with a thiol group (5 0 -GTAAAACGACGACGGCCAGT-SH-3 0 ). A 57-base oligonucleotide (37) included a central sequence complementary to 36 and sides of alternating A and T bases. Oligonucleotide 38 was 50-base poly(dAdT). This oligo-
6.5 Assembly of Biomaterial-Nanoparticle Architectures on Surfaces
Fig. 6-20 A Approach for two-dimensional
assembly of Au nanoparticles using a DNAbased method: (i) a gold nanoparticle is
attached to oligonucleotide 36 through an
Au-thiol interaction; (ii) the oligonucleotide
36 is hybridized with oligonucleotide 37
possessing a complementary base sequence;
(iii) the components are built into a DNA
network consisting of oligonucleotide 38. B
AFM image of the resulting two-dimensional
Au nanoparticle-DNA network. (Adopted from
[154], Figures 1 and 2(a), with permission.)
nucleotide forms a network structure and is hybridized upon contact with 37. All
three oligonucleotides were mixed with Au nanoparticles (5 nm). The resulting
network was adsorbed onto a mica surface, and AFM images were obtained (Figure
6-20B). These images resembled a nanocircuit, but it lacked organization and direct contact between the Au nanoparticles. In order to construct real micro-sized
DNA-templated wires, micro-wires prepared by electrochemical metal deposition
in a porous alumina membrane were bound to thiolated DNA molecules and
reacted with a solid support functionalized with a complementary DNA oligonucleotide [155]. The hybridization process between the surface-confined DNA and
the DNA linked to the microwires yielded ds-DNA bridges binding the microwires
to the surface. It should be noted that in this work the DNA molecules were used
as ‘‘molecular glue’’ to bind micro-sized predesigned wires to the surface instead
of individual nanoparticles, thus providing a significant advance in micro-scale
circuitry.
Another approach to the directed placement of nanoparticles on solid supports by means of DNA is based on the micropatterning of a surface with
DNA molecules. This approach provides the specific binding of complementaryDNA-functionalized nanoparticles at specific domains of the pattern. An aminoderivatized oligonucleotide was deposited onto a chemically modified glass surface
in a form of a pattern by a nanoliter dispensing devise, and was then covalently
bound to the surface [156]. Au nanoparticles (34 nm) functionalized with a complementary oligonucleotide were then reacted with the DNA-patterned surface,
yielding an Au nanoparticle pattern that follows the shape of the DNA pattern. Even
more precise placement of Au nanoparticles was achieved by the ‘‘dip-pen nanolithography’’ that was applied for patterning the primary DNA on the surface [157].
Two kinds of amino-functionalized DNA oligonucleotides, TCTCAACTCGTAA10
(39) and A10 CGCATTCAGGAT (40), were deposited onto a 1,16-mercaptohex-
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
Fig. 6-21 A The use of dip-pen lithography and DNA to
produce a predesigned multinanoparticle pattern. B
Noncontact AFM topography image of such an assembly
consisting of large and small nanoparticles, scale bar is 1 mm.
(Part B reproduced from [157], Figure 1(C), with permission.)
adecanoic acid-functionalized Au surface, and were then covalently bound to it
(Figure 6-21A). The deposition was performed in two sequential steps to yield
a pattern composed of spots of both oligonucleotides. This primary pattern was
reacted with an oligonucleotide of sequence TACGAGTTGAGAATCCTGAATGCG
(41) composed of two domains: one complementary to 39 and the other to 40. The
resulting DNA-functionalized surface provided complementary DNA chains for
two kinds of DNA-functionalized Au nanoparticles: (39)-Au (13 nm) and 40-Au
(31 nm). A pattern was obtained on the DNA-functionalized surface as visualized
by AFM: 39-Au nanoparticles specifically bound to the spots of 40/41 and 40-Au
nanoparticles bound to the spots of 39/41 periodically located on the solid
support (Figure 6-21B).
6.6 Functional Biomaterial-Nanoparticle Structures on Surfaces
Fig. 6-22
The use of DNA as a linker to construct nanoparticle multilayers on surfaces.
Multilayers of nanoparticles can be assembled on solid supports by utilizing
DNA complementarity [158]. For this purpose, a glass surface was functionalized
with a monolayer of an oligonucleotide (42), and then the surface was reacted with
an oligonucleotide (43) composed of two domains, one of which was complementary to 42 while the other provided complemetarity for another oligonucleotide
(44). Oligonucleotide 44-functionalized Au nanoparticles (13 nm) were added to
yield a monolayer of ds-DNA (ds-43/44) attached to Au nanoparticles (Figure 6-22).
Since many of the DNA chains (44) around the Au nanoparticle were not reacted
with the surface-confined DNA (43), they could be reacted with further complementary domains of oligonucleotide 43, thus providing binding sites for 42. A
second kind of Au nanoparticle modified with oligonucleotide 42 was reacted with
this first layer assembly, resulting in a second layer of Au nanoparticles. The repetitive stepwise deposition of Au nanoparticles functionalized with the oligonucleotides 42 and 44 resulted in a multilayer assembly of Au nanoparticles. Each
hybridized Au nanoparticle layer imparted a deeper red color onto the substrate. It
should be noted that this method could allow layer-by-layer deposition of different
(by size or by chemical nature, e.g., Ag, CdS, CdSe) kinds of nanoparticles.
6.6
Functional Biomaterial-Nanoparticle Structures on Surfaces for Sensoric and
Electronic Applications
The unique properties of nanoparticles make nanoparticle-biomaterial conjugates
attractive labels for sensing applications. The optical and electronic sensing of
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
biomaterials on surfaces is a common practice in analytical biochemistry. Thus, the
immobilization of nanoparticle-biomaterial conjugates on surfaces provides a general route for the development of optical or electronic biosensors. Also, the immobilization of metallic or semiconductor nanoparticle-biomaterial conjugates gives a
method for the assembly of electronic circuitry. In such a system, the biomaterial
could provide a structural template for the electronic circuitry. The photonic and
optical properties of the nanoparticles enables imaging of the structural features
of the circuitry, whereas the catalytic functions of the nanoparticles allow their
metallic enlargement to become interconnected with long-range conductivity.
6.6.1
Nanoparticle-Biomaterial Conjugates for Optical Sensing and Analysis
Au nanoparticles linked to bioreceptors provide labeled conjugates that can be used
to follow the biorecognition event at a biosensor surface. Various optical methods
have been employed to detect the association of Au nanoparticles to biochip domains, including, for example the scanometric detection of light scattering, surface
plasmon resonance, resonance-enhanced absorption by nanoparticles, nanoparticle
fluorescence, and enhanced Raman scattering. Although not limited to DNA detection, most of these detection schemes were applied for this purpose.
In a typical setup for scanometric detection, a modified glass slide was mounted
on a microscope stage and illuminated in the plane of the slide with white light.
In this configuration, the slide served as a planar waveguide, preventing any
light from reaching the microscope objective by total internal reflection. Wherever
nanoparticle probes were attached to the surface, however, evanescently coupled
light was scattered out of the slide and was seen as brightly colored spots. This
approach was used for the detection of Au nanoparticle-labeled DNA molecules
specifically bound to a DNA-functionalized surface [159, 160]. Enlargement of the
Au nanoparticles by the reduction of silver salts on them allows a 100-fold amplification of the signal, and thus increases the sensitivity. This method was used to
detect single-base mismatched DNA oligonucleotides immobilized at different domains of the glass support. The high sensitivity required was provided by the silver
deposition, and selectivity was achieved by examining the ‘‘melting’’ profiles of the
spots (the mismatched spot has a lower ‘‘melt’’ temperature due to its lower association constant). Two different DNA sequences could be detected at once when
two Au nanoparticles of different sizes were used to label different oligonucleotides
[161]. Spots with different colors corresponding to two kinds of Au nanoparticles
were detected on the glass support when domains functionalized with different
DNA samples were reacted with the corresponding complementary DNA labeled
with Au nanoparticles 50 nm and 100 nm in diameter. Ag nanoparticles have
similarly been used as antibody labels for the detection of antigens. This setup
provided sensitivity high enough to perform immunoassays based on single-target
detection [162]. Resonance enhancement of the absorptive properties of metal
nanoparticles bound to a surface by biorecognition interactions was also used as an
effective means for biosensor devices [163].
6.6 Functional Biomaterial-Nanoparticle Structures on Surfaces
Au nanoparticles have been used extensively for the detection of bioaffinity
binding by virtue of enhanced surface plasmon resonance (SPR). A dramatic
enhancement of SPR biosensing was achieved for immunosensing [164–166] and
DNA sensing when Au nanoparticles were used [167]. A sensor interface was
modified with antibody units, the surface was reacted with the complementary
antigen component, and finally this affinity assembly was reacted with a secondary antibody labeled with an Au nanoparticle. The association of antigen with an
antibody-functionalized surface can be detected by SPR. However, the change in
the SPR spectrum is larger if the secondary antibody is bound to the surface. It
should be noted that the secondary antibody was able to associate with the modified surface only when the antigen was already bound there, providing an amplification route for the primary recognition event. Amplification is dramatically increased when the secondary antibody is labeled with an Au nanoparticle. The
binding of the Au nanoparticle to the immunosensing interface led to a large shift
in plasmon angle, a broadened plasmon resonance, and an increase in the minimum reflectance, thereby allowing picomolar detection of the antigen. Similarly,
an enhancement in sensitivity by ca. 3 orders of magnitude was obtained in DNA
analysis when Au nanoparticle-functionalized DNA molecules were used as probes
(Figure 6-23) [167]. Surface-enhanced Raman scattering (SERS) of nanoparticlebound substrates allows the amplification of molecular vibrational spectra up to
10 5 -fold [69]. Although the technique has not been widely applied for detecting
Fig. 6-23 Surface plasmon resonance (SPR)
spectroscopy use for DNA analysis enhanced
by Au nanoparticles: A SPR spectrum of the
DNA-primer-functionalized surface, B SPR
spectrum after hybridization with the DNA-
analyte molecules, and C SPR spectrum after
hybridization with the DNA-analyte molecules
conjugated with Au nanoparticles. (Adopted
from [167], Figure 1, with permission.)
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
biorecognition events on surfaces, it was demonstrated that a cytochrome c-Au
nanoparticle conjugate associated with an Ag surface revealed an SERS spectrum
[168, 169]. Semiconductive nanoparticles (e.g., CdS, CdSe) can be used as fluorescent labels for immunosensing and DNA sensing, providing tunable wavelength,
narrow emission peaks, and 100-fold higher stability than molecular fluorescent
dyes [69].
6.6.2
Electronic Transduction of Biosensing Events by Nanoparticle-BiomaterialFunctionalized Arrays
Redox-enzymes can be contacted with electrodes by the tethering of redox units to
the proteins [170–173] or by the surface reconstitution of apo-proteins on electronrelay cofactor units [174–178]. Recently, the electrical contacting of redox-enzymes
was accomplished by the reconstitution of the apo-flavoenzyme apo-glucose
oxidase (apo-GOx) with a 1.4 nm Au55 -nanoparticle functionalized with N 6 -(2aminoethyl)-flavin adenine dinucleotide (FAD cofactor amino-derivative, 45). The
conjugate produced was assembled on a thiolated monolayer using various dithiols
(46–48) as linkers (Figure 6-24A) [179]. Alternatively, the FAD-functionalized Au
nanoparticle could be assembled on a thiolated monolayer associated with an electrode, with apo-GOx subsequently reconstituted on the functional nanoparticles
(Figure 6-24B). The enzyme-electrodes prepared by these two routes reveal similar protein surface coverages of ca. 1 1012 molecm2 . The nanoparticlereconstituted glucose oxidase layer is electrically contacted with the electrode without any additional mediators, and it stimulates the bioelectrocatalyzed oxidation of
glucose (Figure 6-24C). The resulting nanoparticle-reconstituted enzyme electrodes
reveal unprecedentedly efficient electrical communication with the electrode (electron transfer turnover rate ca. 5000 s1 ). This electrical contacting makes the
enzyme-electrode insensitive to oxygen or to common oxidizable interferants such
as ascorbic acid.
A related electrochemical method was employed for the Au nanoparticle-based
quantitative detection of the 406-base human cytomegalovirus DNA sequence
(HCMV DNA) [180]. The HCMV DNA was immobilized on a microwell surface
and hybridized with an oligonucleotide-modified Au nanoparticle. The resulting
surface-immobilized Au nanoparticle double-stranded assembly was treated with
HBr/Br2, resulting in the oxidative dissolution of the gold particles. The solubilized
Au 3þ ions were electrochemically reduced and accumulated, and then determined
by anodic stripping voltammetry using a sandwich-type screen-printed microband
electrode (SPMBE). The combination of the sensitive detection of Au 3þ ions at the
SPMBE by non-linear mass transfer of the ions and the release of a large number
of Au 3þ ions upon the dissolution of the particle associated with a single recognition event provides an amplification path that enables the detection of the HCMV
DNA at a concentration of 5 1012 M.
The conductivity stimulated by patterned nanoparticle bridges provides an alternative route for designing miniaturized biosensors. This was exemplified by the
6.6 Functional Biomaterial-Nanoparticle Structures on Surfaces
Fig. 6-24 Electrical ‘‘wiring’’ of glucose
oxidase (GOx) by apo-enzyme reconstitution
with an Au nanoparticle functionalized with a
single FAD cofactor unit. A Reconstitution
process performed in solution followed by the
assembly adsorption onto a dithiol-modified
Au electrode. B Au-FAD conjugate adsorption
onto a dithiol-modified Au electrode followed
by the reconstitution of the apo-GOx at the
interface. C Calibration plot of the electrocatalytic current developed by the reconstituted
GOx electrode in the presence of different
concentrations of glucose.
design of a miniaturized immunosensor based on Au nanoparticles and their
catalytic properties (Figure 6-25) [181]. Latex particles stabilized by an anionic protective layer are attracted to a gap between micron-sized Au electrodes by the
application of a non-uniform alternating electric field between the electrodes
(dielectrophoresis). Removal of the protective layer from the latex particles by an
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
Immunosensing at micro-sized Au electrodes based
on the change of conductivity between the Au strips upon
binding of Au nanoparticles followed by silver deposition.
Fig. 6-25
oppositely charged polyelectrolyte results in the aggregation of the latex particles
and their fixation in the gap. Adsorption of protein A on the latex surface yields a
sensing interface for the human immunoglobulin (IgG) antigen. The association
of human immunoglobulin on the surface is probed by the association of the secondary Au-labeled anti-human IgG antibodies to the surface, followed by the catalytic deposition of a silver layer on the Au particles. The silver layer bridges the gap
between two microelectrodes, resulting in a conductive ‘‘wire’’. Typical resistances
between the microelectrodes were 50–70 W, whereas control experiments generated resistances >10 3 W. The method enabled the analysis of human IgG with a
detection limit of ca. 2 1013 M.
DNA detection has also been performed using microelectrodes fabricated on a
silicon chip (Figure 6-26) [182]. A probe nucleic acid (49) was immobilized on the
SiO2 interface in the gap separating the microelectrodes. The target 27-mer
nucleotide (50) was then hybridized with the probe interface, and subsequently
a nucleic acid (51)-functionalized Au nanoparticle was hybridized with the free
3 0 -end of the target DNA. The Au nanoparticle catalyzed hydroquinone-mediated
reduction of Agþ ions, resulting in the deposition of silver on the particles, lowering the resistance between the electrodes. Single-base mutants of the analyte
oligonucleotide (50) were washed off from the capture-nucleic acid (49) by the
application of a buffer of an appropriate ionic strength, so were not detected. A
difference of 10 6 in the gap resistance was found between the analyte and the
6.6 Functional Biomaterial-Nanoparticle Structures on Surfaces
Fig. 6-26 The use of a DNA-nanoparticle conjugate and
subsequent silver deposition to connect two microelectrodes,
as a means to sense a DNA analyte.
mutants. The low resistances between the microelectrodes were found to be controlled by the concentration of the target DNA, and the detection limit for analysis
is in the region of 5 1013 M. This sensitivity translates to ca. 1 mgmL1 of
human genomic DNA or ca. 0.3 ngmL1 of DNA from a small bacterium. These
concentrations suggest that the DNA may be analyzed with no pre-PCR amplification.
Photoelectrochemical transduction of DNA recognition processes has been
shown by use of semiconductor (CdS) nanoparticles modified with nucleic acids
[183]. Semiconductor CdS nanoparticles (2:6 G 0:4 nm) were functionalized with
the two thiolated nucleic acids 52 and 53 that are complementary to the 5 0 and 3 0
ends of a target DNA (54). An array of CdS nanoparticle layers was constructed on
an Au electrode by a layer-by-layer hybridization process (Figure 6-27). Illumination of the array resulted in the generation of a photocurrent. The photocurrent
increases with the number of CdS nanoparticle layers, and the photocurrent action
spectra follow the absorbance features of the CdS nanoparticles, implying that the
photocurrents originate from the photoexcitation of the CdS nanoparticles. That is,
photoexcitation of the semiconductor causes the formation of an electron-hole pair.
Transfer of conduction band electrons, and the concomitant transfer of electrons
from a sacrificial electron donor to the valence band holes, yields a steady-state
photocurrent in the system. The ejection of the conduction-band electrons into the
electrode occurred from nanoparticles in intimate contact with the electrode support. This is supported by the fact that Ru(NH3 )3 6þ units (E 0 ¼ 0:16 V vs SCE)
that are electrostatically bound to the DNA enhance the photocurrent from the
DNA-CdS array. That is, the Ru(NH3 )6 3þ units act as an electron wiring element
that facilitates electron hopping to the electrode through the DNA tether from CdS
particles that lack contact with the electrode. The system is important not only as it
demonstrates the use of photoelectrochemistry as a transduction method for DNA
sensing, but also since the system reveals the nano-engineering of organized DNA-
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
The construction of CdS-nanoparticle/DNA
superstructures and their use for the generation of
photocurrents.
Fig. 6-27
tethered semiconductor nanoparticles on conductive supports. These latter nanoengineered structures are the first step toward electronic nano-circuitry (see discussion in Section 6.6.3).
Nanoparticles provide a high-mass method of labeling nucleic acid-DNA, making microgravimetric sensing methods (e.g., QCM) ideal for the detection of recognition events. Furthermore, as nanoparticles act as catalysts for the deposition of
metals, even higher mass changes may be stimulated. For a quartz piezoelectric
crystal (AT-cut), the crystal resonance frequency changes by Df when a mass
change Dm occurs on the crystal according to Eq. (1) (the Sauerbrey equation),
where f0 is the fundamental frequency of the quartz crystal, Dm is the mass
change, A is the piezoelectrically active area, rq is the density of quartz (2.648
gcm3 ), and mq is the shear modulus (2:947 10 11 dyncm2 for AT-cut quartz).
Df ¼ 2 f 0 2
Dm
A ðmq rq Þ 1=2
ð1Þ
Microgravimetric DNA detection using nucleic acid-functionalized Au nanoparticles as ‘‘nano-weights’’ was accomplished by the hybridization of a target
DNA (55) (50 -AGCGTAGGATAGATATACGGTTCGCGC-30 ) to an Au-quartz crystal
modified with a probe oligonucleotide, followed by the hybridization of the interface with an Au nanoparticle functionalized with a nucleic acid (56) (50 -HS-(CH2 )6 GCGCGAACCGTATA-30 complementary to the free 3 0 -segment of the target DNA
(55) [184, 185]. Further amplification of the response was reported by the use of a
secondary Au nanoparticle that is functionalized with the nucleic acid (57) (50 TCTATCCTACGCT-(CH2 )-SH-30 ) that is complementary to the 5 0 -segment of the
6.6 Functional Biomaterial-Nanoparticle Structures on Surfaces
target DNA (55). Hybridization of the secondary nucleic acid-labeled Au nanoparticle with 55 followed by hybridization to the primary nanoparticle layer, yields a
‘‘second generation’’ of Au nanoparticles reminiscent of the growth of dendrimers
[185, 186]. Concentrations as low as 1 1010 M of DNA could be sensed by the
amplification of the target DNA by the nucleic acid-functionalized Au nanoparticle
labels.
A different approach to the amplified quartz-crystal-microbalance analysis of
DNA utilizes the catalytic metal deposition on the nanoparticle labels [187]. Figure
6-28A depicts the amplified detection of the 7229-base M13f DNA using the catalytic deposition of gold on an Au nanoparticle conjugate [188]. The primer (58) is
assembled on an Au/quartz crystal. After hybridization with M13f DNA (59), the
Fig. 6-28 A Amplified detection of the 7229-base M13f DNA
using the catalytic deposition of gold on an Au nanoparticle
conjugate. B Analysis of a single base mismatch in DNA using
the catalytic deposition of gold on an Au nanoparticle
conjugate.
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
double-stranded assembly is replicated in the presence of dATP, dGTP, dUTP and
biotinylated-dCTP (B-dCT). The resulting biotin-labeled replica is then reacted with
a streptavidin-Au nanoparticle conjugate (60), and the resulting Au-labeled replica
is subjected to the Au nanoparticle-catalyzed deposition of gold by the NH2 OHstimulated reduction of AuCl4 . The replication process represents the primary
amplification, since it increases the mass associated with the crystal and simultaneously generates a high number of biotin labels for the association of the
streptavidin-Au nanoparticle conjugate. The binding of the nanoparticle conjugate
represents the secondary amplification step for the analysis of the M13f DNA. The
third step, catalyzed precipitation of the metal, gives the highest amplification. The
M13f DNA could be sensed by this method with a lower detection limit of ca.
1 1015 M. This amplification route was also applied for the analysis of a single
base mismatch in DNA, exemplified by the analysis of the mutant (61a) that
includes the single base substitution of the A-base in the normal gene (61) with a
G-base (Figure 6-28B). The analysis of the mutant is performed by the immobilization of probe DNA (62) that is complementary to the normal gene (61) as well as
to the mutant (61a) up to one base prior to the mutation site, on the Au-quartz
crystal. Hybridization of the normal gene or the mutant with this probe interface,
followed by the reaction of the hybridized surfaces with biotinylated dCTP, BdCTP, in the presence of Klenow fragments, incorporates the biotin-labeled base
only into the assembly that includes the mutant. The subsequent association of the
streptavidin-Au conjugate (60), followed by the catalyzed deposition of gold on the
Au nanoparticles, amplifies the analysis of the single base mismatch in 61a. Figure
6-29, curve (a), shows the microgravimetric detection of the mutant 61a by this
method, where the normal gene (61) does not alter the frequency of the crystal
Microgravimetric detection of a
single-base mutant enhanced by the catalytic
deposition of gold on an Au nanoparticle
conjugate. A frequency change is observed
with a mutant (a) and with a normal DNA
Fig. 6-29
sequence (b). Arrow (1) shows the attachment
of the Sav-Au conjugate. Arrow (2) shows
catalytic deposition of gold on the Au
nanoparticles. (Adopted from [188], Scheme 1
and Figure 1(B), with permission.)
6.6 Functional Biomaterial-Nanoparticle Structures on Surfaces
(Figure 6-29, curve (b)). Using this method, the mutant could be detected with a
detection limit of 5 1013 M.
The immobilization of nanoparticles on surfaces may also be used to yield high
surface area electrodes [189]. Enhanced electrochemical detection of nucleic acids
was reported by the roughening of flat gold electrodes with an Au nanoparticle
monolayer [190]. The roughening of an Au-quartz crystal with a monolayer consisting of Au nanoparticles has also been employed for the enhanced microgravimetric analysis of DNA [191].
6.6.3
Biomaterial-Nanoparticle Arrays for Electronic Circuitry
As discussed above (cf. Chapter 6.5.2), DNA molecules can be loaded with metallic
or semiconductive nanoparticles, which form wires or networks upon their deposition onto a solid. In order to use these nanowires for electronic circuitry, they
should be electrically connected to the external world, and they need to be conductive. Since nanoparticles are loaded onto the DNA template with gaps between
them, the issue of electrical conductivity is important. Electron transport through
DNA has been one of the most intensively debated subjects in chemistry over the
past ten years [192, 193], and it is still under extensive theoretical and experimental
investigation. In order to examine the electrical conductivity of ds-DNA spacers
connecting Au nanoparticles, thiol-derivatized oligonucleotides were adsorbed on
Au nanoparticles, and then the DNA-functionalized Au nanoparticles were bridged
with DNA chains composed of double-stranded helices of various lengths (24, 48
or 72 bases). These helices were terminated on both sides with single-stranded
domains complementary to the oligonucleotides bound to the Au nanoparticles
[194]. The resulting Au nanoparticle aggregates linked with ds-DNA spacers were
deposited on an electrically non-conductive solid support, and their conductivity
was measured by the four-probe method. Surprisingly, the conductivities of the
aggregates formed from all three linkers ranged from 105 to 104 S cm1 at room
temperature, and they showed similar temperature-dependent behavior. The similarity of the electrical properties of the aggregates originates from the fact that the
DNA spacers are compressed, thus providing small and similar distances between
the Au nanoparticles. Accordingly, the measured conductivity parameters reflect
the electrical properties of the metallic nanoparticles separated by short gaps.
The conductivity of metallic nanoparticle aggregates can be enhanced upon the
chemical deposition of another metal (e.g., Ag deposition on Au aggregates). DNAfunctionalized Au nanoparticles were placed between microelectrodes by hybridization with complementary DNA molecules immobilized in the non-conductive
gap between two electrodes [195]. Then the nanoparticles were subjected to the
reductive deposition of Ag to enlarge the metallic species and bridge the gaps between them with deposited Ag. The subsequent value of conductivity was found
to be dependent on the original loading of the gap with Au nanoparticles, thus
allowing the quantitative measurements of the DNA-functionalized Au nanoparticles bound to the surface.
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
Fig. 6-30 A The construction of a wire bridging two
microelectrodes by the deposition of nanoparticles on a
bridging DNA strand, followed by silver deposition. B I–V
curves obtained from the structures produced. (Part B adopted
from [196], Figure 6-4(a), with permission.)
A single nanowire produced on a DNA template bridging two micro-size electrodes was constructed [196, 197]. Two micro-electrodes facing each other (12–16
mm separation) were functionalized with 12-base oligonucleotides that were then
bridged with a 16 mm long l-DNA (Figure 6-30A). The resulting phosphate units of
the DNA bridge were loaded with Agþ ions by ion-exchange, which were then reduced to Ag metal by hydroquinone. The small Ag aggregates produced along the
DNA backbone were then used as catalysts for further reductive deposition of silver, eventually leading to the formation of an Ag nanowire. This micrometer-sized
element had a typical width of 100 nm and a granular morphology (as determined
by AFM). Electrical measurements revealed non-linear I–V curves (Figure 6-30B).
A similar approach was used to generate highly conductive nanowires bridging
macroscopic Au electrodes [198]. The DNA molecules were positioned between
macroscopic Au electrodes and loaded with Pd 2þ ions. The ions were then chemi-
6.7 Biomaterial-Functionalized Magnetic Particles
cally reduced to continuous metallic wires, providing specific conductivity only one
order of magnitude below than of bulk palladium.
6.7
Biomaterial-Functionalized Magnetic Particles
Magnetic particles (microspheres, nanospheres, and ferrofluids) are widely studied
and applied in various fields of biology and medicine such as magnetic targeting
(drugs, genes, radiopharmaceuticals), magnetic resonance imaging, diagnostics,
immunoassays, RNA and DNA purification, and cell separation and purification
[199]. These magnetic particles are generally of the ‘‘core-shell’’ type: biological
species (cells, nucleic acids, proteins) are connected to the magnetic ‘‘core’’ through
organic linkers, often organized as a polymeric ‘‘shell’’ around the core. The magnetic core is usually composed of Fe3 O4 , and its primary modification with an
organic ‘‘shell’’ can include the adsorption of an organic polymer or the covalent
attachment of organosilane molecules. Organic functional groups introduced by
these techniques allow the coupling of biomolecules to the organic ‘‘shell’’. Several
synthetic approaches have been applied to couple artificial magnetic particles to
biomolecules, which were then used for various bioanalytical applications. For example, antibody molecules were adsorbed on synthetic Fe3 O4 magnetic particles
and were then used for specific binding to cells and the separation of the cells in
an external magnetic field [200, 201]. DNA molecules were reacted with a mixture
of Fe 2þ and Fe 3þ ions that were electrostatically associated with DNA chains [202].
These iron ions were then chemically reacted to give Fe3 O4 magnetic particles
associated with the DNA molecules. The labeled DNA could hybridize with complementary oligonucleotides, and the magnetic particles linked to the DNA molecules allowed the separation of the labeled from the non-labeled DNA strands.
The natural protein ferritin can be used to generate precisely sized and shaped
magnetic particles in a protein cavity [203]. For this purpose, the non-magnetic
natural core composed of 5 Fe 2 O3 9 H2 O was chemically removed from the protein cavity, and an artificial magnetic core composed of magnetite (Fe3 O4 ) [104] or
magnetite/maghemite (Fe3 O4/g-Fe 2 O3 ) [204] was generated inside it. The natural
protein shell remains, providing binding groups for the covalent attachment of
biomolecules. Some bacteria (e.g., Magnetospirillum magnetotacticum MS-1 [205],
Magnetospirillum sp. MGT-1 [206], Magnetospirillum sp. AMB-1 [207], and Magnetospirillum gryphiswaldense [208]) produce natural magnetic particles (Fe3 O4 , 50–
100 mm in size) aligned in chains and enveloped by a lipid membrane [209, 210].
These naturally produced lipid-covered magnetic particles can be isolated from the
parent bacteria [211] and modified with biomolecules using bifunctional coupling
reagents (e.g., glutaric dialdehyde). They have been applied for fluoroimmunoassay
[212–214], chemiluminescence immunoassay [215], mRNA recovery [216], and as
DNA carriers [216, 217]. Most of these techniques are based on the magnetic separation of the magnetic nanoparticle-linked biomolecules from the non-labeled
biomolecules.
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
The immobilization of enzymes on magnetic particles yields biocatalytically
active particles. In one example, alcohol dehydrogenase was covalently bound to
Fe3 O4 magnetic particles, and the immobilized enzyme continued to demonstrate
high biocatalytic activity [218, 219].
Most of the applications of biomaterial-magnetic particle hybrid conjugates
involve the concentration, separation, regeneration, mechanical translocation, and
targeting of biomaterials. As these subjects do not deal directly with nanoparticlebiomaterial conjugates, we will not cover them here. Some recent research,
however, has been directed to the coupling of functional magnetic particles with
biomaterial to yield magnetically-controlled functions. Magnetic particles were
modified with redox-relay groups such as ferrocene or bipyridinium units and applied for the magneto-switched activation or deactivation of the bioelectrocatalytic
functions of redox-enzymes (Figure 6-31A) [220]. Ferrocene (63)-functionalized
magnetite particles were used for the magnetic switching of the bioelectrocatalyzed
oxidation of glucose in the presence of glucose oxidase (GOx) (Figure 6-31B).
Magnetic attraction of the particles to the electrode enabled the oxidation of the
ferrocene-linked redox units that mediated the GOx-biocatalyzed oxidation of glucose. Positioning of an external magnet above the electrochemical cells lifts the
magnetic particles from the electrode, and the bioelectrocatalyzed oxidation of glucose is switched off. Figure 6-31C shows the magneto-switched bioelectrocatalyzed
oxidation of glucose by positioning the external magnet below and above the electrode. Similar magnetic switching of the bioelectrocatalyzed reduction of nitrate
in the presence of nitrate reducatase (NR) was accomplished by the application
of N,N 0 -bipyridinium (64)-functionalized magnetic particles. Similarly, magnetic
nanoparticles linked to nicotinamide adenine dinucleotide through PQQ were
used for the magnetic switching of the bioelectrocatalyzed oxidation of lactate to
pyruvate using lactate dehydrogenase (LDH) [221]. These particles were employed
for the dual selective sensing of lactate and glucose in the presence of LDH and
GOx [222] (Figure 6-32). An Au electrode was functionalized with a ferrocenelabeled monolayer, and PQQ-NADþ -modified particles were suspended in electrolyte solution. Removal of the magnetic particles from the electrode by means of
the external magnet enables the electrochemical detection of glucose at potentials
more positive than 0.32 V (vs SCE) by the ferrocene-mediated bioelectrocatalyzed
oxidation of glucose in the presence of GOx (Figure 6-33A). Attraction of the functional magnetite particles to the electrodes by the external magnet enables the
amperometric analysis of lactate in the potential range 0.13 V < E < 0.32 V
(Figure 6-33B).
A novel approach to DNA analysis [223] (Figure 6-34) involved the use of
magnetic particles modified by a nucleic acid (65) that hybridizes with a biotinylated target DNA (66). The resulting hybridized complex was then reacted with
a streptavidin-Au nanoparticle conjugate (67). A silver shell was catalytically generated on the Au nanoparticles linked to the DNA molecules, and then the magnetically labeled assembly was moved to an electrode by means of an external
magnet, where electrochemical stripping of the Ag layer provided an amperometric
signal for the quantitative determination of the analyte DNA.
6.7 Biomaterial-Functionalized Magnetic Particles
Fig. 6-31 A Synthesis of redox-relay functionalized magnetic particles. B Magnetoswitched bioelectrocatalyzed oxidation of
glucose in the presence of glucose oxidase and
ferrocene-functionalized magnetic particles.
C Cyclic voltammograms corresponding to the
magneto-switched bioelectrocatalyzed oxidation of glucose: (a) Ferrocene-functionalized
magnetic particles are attracted to the
electrode by means of an external magnet.
(b) Ferrocene-functionalized magnetic particles
are retracted from the electrode by means of an
external magnet. Inset: Cyclic switching ‘‘ON’’
and ‘‘OFF’’ the bioelectrocatalyzed oxidation of
glucose by the reversible attraction and retraction of the functionalized-magnetic particles to
and from the electrode, respectively. (Adopted
from [220], Schemes 1 and 2, Figure 1(B), with
permission.)
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
Magneto-switched dual biosensing of glucose and
lactate. (Adopted from [222], Figure 1, with permission.)
Fig. 6-32
Fig. 6-33 A Cyclic voltammograms of the
ferrocene-monolayer-functionalized Au electrode when the NADþ /PQQ-functionalized
magnetic particles are retracted from the electrode by the external magnet: (a) in the presence of GOx, 1 mg mL1 ; LDH, 2 mg mL1 ;
(b) in the presence of GOx, 1 mg mL1 ; LDH,
2 mg mL1 ; glucose, 50 mM; lactate, 20 mM.
B Cyclic voltammograms of the ferrocene-
monolayer-functionalized Au electrode in the
presence of NADþ /PQQ-functionalized
magnetic particles attracted to the electrode
by the magnet: (a) in the presence of GOx,
1 mg mL1 ; LDH, 2 mg mL1 ; (b) in the
presence of GOx, 1 mg mL1 ; LDH, 2
mg mL1 ; glucose, 50 mM; lactate, 20 mM.
(Adopted from ref. 222, Figure 6-3, with the
permission.)
6.7 Biomaterial-Functionalized Magnetic Particles
Fig. 6-34 Assembly of DNA molecules at magnetic particles
followed by their association with Au nanoparticles and further
silver deposition, and the DNA analysis by electrochemical Ag
stripping.
The mechanical rotation of magnetic particles using an external rotating magnet
has also been used as an amplification method for electrochemical sensing [224].
In these systems, chemical transformations are controlled by convection of the
solution and its constituents rather than by diffusion. As a result, mass transport
to the magnetic particles is controlled by the rotational speed of the particles
around the electrode axis. Pyrroloquinoline quinone (68)-functionalized magnetite particles were employed for the electrocatalyzed oxidation of NADH under
rotation (Figure 6-35A). The electrocatalytic anodic currents are controlled by the
circular rotation speed of the particles (Figure 6-35B). Similarly, the ferrocene (69)functionalized particles were used to control the bioelectrocatalyzed oxidation of
glucose in the presence of GOx by means of the circular rotation of the particles
produced by the external magnetic rotor (Figure 6-36A). The bioelectrocatalytic
current provided by the soluble GOx and ferrocene mediator linked to the rotating
magnetic particles were dependent on the particle rotation speed (Figure 6-36B).
The calibration plots derived from the cyclic voltammograms show a dramatic
influence of the particle rotation speed on the sensitivity of the bioelectrocatalytic
system to glucose (Figure 6-36C). This method represents a novel route to the
413
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6 Biomaterial-Nanoparticle Hybrid Systems: Synthesis, Properties, and Applications
Fig. 6-35 A Electrocatalytic oxidation of NADH
enhanced by circular rotation of the PQQfunctionalized magnetic particles induced by
the rotation of an external magnet. B Cyclic
voltammograms corresponding to the electrocatalytic oxidation of NADH upon different
rotation speeds of the PQQ-functionalized
magnetic particles.
amplification of bioelectronic detection schemes or the enhancement of biological
recognition processes. The labeling of biological components by magnetic particles
(e.g., antibodies or nucleic acids), and their external rotation by means of the
magnetic rotor, may enhance the formation of antigen-antibody complexes or
hybridization, respectively, by the agitation of the solution. Similarly, biorecognition events transduced by means of bioelectrocatalytic labels may be amplified by
coupling the bioelectrocatalytic label to magnetic labels and rotating the conjugated
biorecognition complexes with an external magnetic rotor.
6.7 Biomaterial-Functionalized Magnetic Particles
Fig. 6-36 A Electrocatalytic oxidation of
glucose in the presence of glucose oxidase
enhanced by circular rotation of the ferrocenefunctionalized magnetic particles induced by
the rotation of an external magnet. B Cyclic
voltammograms corresponding to the
electrocatalytic oxidation of glucose upon
different rotation speeds of the ferrocenefunctionalized magnetic particles. C Calibration
plots obtained with different rotation speeds.
Acknowledgment
Our research on biomaterial-nanoparticle hybrid systems is supported by the
German-Israeli Foundation (GIF).
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7
Conclusions and Perspectives
All Authors
Nanoparticles, whether they originate from semiconductors or from metals, are
generally associated with the appearance of novel properties. The most far-reaching
consequence of their small size is without doubt the change in their electronic
properties. Next in importance are changes in other properties, both physical and
chemical. These changes are what ultimately constitute the value of nanosciences
and nanotechnology. Novel properties become available without making new materials!
If particles are small enough, they become electronically comparable to atoms
and molecules, following quantum mechanical rules instead of the laws of classical
physics. That is why nanoparticles are called ‘‘quantum dots’’ or sometimes ‘‘artificial atoms’’. The reason for this behavior is the disappearance of band structures
and the formation of discrete energy levels, these being linked to an increase in the
already existing band gaps in the case of semiconductor particles and the formation of quantized states in the case of metal nanoparticles. Two-dimensional
arrays of quantum dots as well as one-dimensional arrangements (quantum wires)
therefore reach a new dimension of quality with respect to possible future applications. The synthesis of all types of quantum dots is therefore a challenge for both
chemists and physicists. Their easy availability is of decisive importance if nanoparticles are to play an industrial role, and in fact II-VI, III-V as well as Ib-VI
semiconductor nanoparticles are routinely accessible by various wet-chemical
methods in gram amounts.
Even the most difficult step in all synthetic procedures, the achievement of
monodispersity of the particles, can be controlled satisfactorily. This is true of both
semiconductor and metal nanoparticles. Because of the long history of metal
nanoparticles, there has been plenty of time to develop perfect syntheses. One
decisive finding opened up completely novel strategies for both semiconductor and
metal nanoparticles: the use of protecting ligand shells, usually consisting of organic molecules. Amines, phosphines, phosphine oxides, thiols, and numerous
other classes of compounds have been used. Such coordinating molecules not only
support the synthesis of monodisperse nanoparticles by their kinetic and sizeNanoparticles: From Theory to Application. Edited by Günter Schmid
Copyright 8 2004 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
ISBN: 3-527-30507-6
7 Conclusions and Perspectives
limiting functions; they also make the particles soluble in solvents, depending on
the chemical nature of the ligand molecules. Here, traditional chemical procedures
such as chromatography, crystallization, etc. can be used for purification. Semiconductor nanoparticles include the II–VI type, e.g., CdS, CdSe and CdTe, the III–
V type, mainly represented by GaN, GaP, GaAs or InP and InAs, and the Ib–VI
type, available as copper selenide and telluride nanoclusters and as silver and gold
selenide and telluride species. However, metal nanoparticles have also had a remarkable renaissance after a hundred years’ sleep. This is especially true of magnetic nanoparticles. Apart from leading to a general increase in our knowledge of
magnetism, these offer unimagined application possibilities.
Although magnetic nanoparticles are currently being used in several industrial
and medical applications, systematic research into the magnetic properties of nanoscale particles is still lacking. The reason for this is clearly linked to a lack of highquality samples. Against this background, exciting fundamental studies on the size
dependence of the magnetic properties in dilute systems can be expected. Questions concerning magnetic phase transitions such as the transition from the ferromagnetic to the superparamagnetic state will attract considerable attention. Going
a step further, one may expect fundamentally new results for highly ordered dense
magnetic systems. The alignment of the spins of neighboring particles relative to
each other will be especially interesting for magnetic transport properties such as
giant magnetic resistance. The possibility of preparing perfectly ordered domains
of almost single-sized and well-defined particles extending over wide areas will lead
to a major step forward in the exciting field of spintronics. Exciting developments
in the use of magnetic nanoparticles in biology and medicine are also expected.
Smart magnetic nanoparticles used in an improved hyperthermal therapy in which
tumor cells are recognized and well-defined and controllable magnetic labels for
the investigation of the nanostructure and dynamics of cells are two goals for the
future.
The fascinating electronic ‘‘inner life’’ of nanoparticles is only of academic relevance if we fail to use them in nanoelectronics, optoelectronics, and other sophisticated techniques. One precondition for such use is that we must learn how to
organize nanoparticles in three dimensions, but even more importantly in one and
two dimensions. Developments during the last decade have resulted in impressive
progress. Self-assembly techniques in particular have been so much improved that
nowadays semiconductor as well as metal nanoparticles can be arranged in perfectly ordered assemblies of remarkable extensions. Our first ideas about interparticle relationships became known from these 2D arrangements. As it turned
out, this kind of ‘‘chemistry’’ between the particles, among other things, determines the nature of their collective behavior.
It is rather more difficult to routinely organize nanoparticles in one direction. To
do this, appropriate templates are indispensable. Several approaches using polymer molecules or DNA look interesting, but still do not give results that could be
regarded as a breakthrough.
On the other hand, techniques have been developed that even allow writing with
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424
7 Conclusions and Perspectives
nanoparticles. Dip-pen lithography is one of these methods, accompanied by
‘‘nano-electrochemical’’ modifications of surfaces, enabling the nanoparticles to be
fixed in a final step.
Much progress has been achieved in understanding charge transport phenomena through nanostructured materials. The electrical properties of nanoparticles
are mainly determined by the Coulomb charging energy, and the phenomena can,
in principle, be understood in terms of single electron tunneling (SET) within the
framework of the ‘‘orthodox theory of SET’’. Furthermore, in the smallest particles,
like Au55 , additional effects due to the discrete nature of the energy states occur,
and need to be taken into account in the discussion of the electrical transport
phenomena and the characteristic excitation energies. This makes the physics of
these structures much more complicated. At least, the ‘‘orthodox theory’’ has to be
extended for this size regime, and since the electron distribution is then no longer
a pure Fermi distribution, electrical capacitances cannot be discussed classically,
and even the tunneling time for electrons can be affected.
These effects have not been studied experimentally or theoretically in much
detail. Thus, most of the presently applied models are at least incomplete with
respect to the huge variety of different control parameters in the chemical design
of molecularly organized nanoparticles. Nevertheless, applying chemically tailored
metal nanoparticles in SET might be an alternative approach to the main-stream
developments, like CMOS technologies, in the future. But it remains an open
question whether these nanoparticles or other molecular objects, like nanotubes or
organic/organometallic molecules, remain stable under the operating conditions of
a high-performance electrical circuits. With respect to circuit design, precise control over size and topology is needed. At the same time, one needs to keep in mind
the fact that a 10 nm CMOS device, the technological fabrication of which is one
of the great challenges of the next decade, consists of a few thousand atoms and
consequently contains only 5–10 doping atoms in the relevant volume fraction of
the respective semiconductor. It will be an extraordinarily difficult task to find
methods which allow precise control over the distribution of the doping sites over
the respective device elements. In contrast, a chemically designed molecule is precise on the atomic level and available on a molar scale. Furthermore, SET devices,
as they can be built up with chemically tailored building blocks, control the motion
of just a few electrons instead of several hundreds or thousands of them. Since
power consumption is roughly proportional to the number of electrons transferred
from a voltage source to the ground in logic operations, SET devices could reduce
the power consumption drastically.
It is evident, for any kind of chemical approach, that reliance on the present-day
paradigm of information exchange in logic elements or memory devices will require defect-tolerant computer architectures with suitable, intelligent software. On
the other hand, it may be possible to develop alternative routes which require more
complex architectures and different principles of information exchange where the
problems associated with redundancy, current density, performance speed, and
especially contact resistance and charge redistribution at the lead-to-molecule in-
7 Conclusions and Perspectives
terface are less relevant. Solving these problems will be one of the great challenges
in chemistry.
The fact that nanoparticles and biomaterials (such as enzymes, antibodies, and
nucleic acids) are of similar dimensions turns hybrid systems, consisting of both,
into attractive nanoelements or building blocks of nanostructures and devices.
Several functions and applications of biomaterial-nanoparticle hybrid systems,
including analytical applications, signal-triggered electronic functions, nanostructures for circuitry, and the assembly of devices, are of importance. Some of these
functions represent viable technologies, whereas others are still in an embryonic
phase that requires additional fundamental research. The analytical applications of
nanoparticle-biomaterial systems have advanced tremendously in the last decade.
The understanding of the unique optical properties of nanoparticles and the
photophysics of coupled interparticle interactions has enabled these particles to be
used as optical labels for recognition events. Similarly, the catalytic properties of
biomaterials have enabled biorecognition events to be amplified. If nanoparticlebiomaterial hybrid systems are assembled on surfaces, the electronic detection of
biorecognition events can become feasible by means of electrical conductivity or
piezoelectric or photoelectrochemical transduction.
Novel techniques, such as electronic sensing amplified by the movement of
magnetic particles by an external magnetic motor, represent a new analytical concept. Many future applications of rotating magnetic particles are envisaged in the
detection of antigen-antibody, nucleic acid-DNA, protein-DNA, or receptor-cell interactions. Also, enhanced reactivity at rotating magnetic particles (e.g., accelerated
hybridization or polymerization) may facilitate analytical processes. Future analytical perspectives of nanoparticle-hybrid systems may involve the mechanical actuation of microelements. The detection of biomaterials by the mechanical movement
of microlevers due to surface stress interactions is well established. The integration
of charged nanoparticles as labels for such biorecognition events may enhance the
surface stress and provide mechanical activation of micro-objects.
The electronic triggering of redox-proteins by the incorporation of nanoparticles
represents a novel strategy for the electrical contacting of redox-enzymes with their
macroscopic environment. The use of other nano-objects to facilitate electrical
contacting, such as the application of metallic nano-rods or carbon nanotubes, is
a future challenge. The magnetic triggering of redox-proteins will certainly find
applications in the tailoring of complex sensor systems and the development of
new surface-patterning methods.
The biggest challenges are ahead of us, however, and exciting systems are
envisaged for the future. Besides the interesting nano-architectures that may be
generated, fundamental problems, such as probing charge transport phenomena
in nanoparticle-biomaterial systems, may be highlighted. Nonetheless, the organization of nanoparticle architectures on biomaterial templates involves relatively
simple motifs that do not make full use of the biomaterial’s catalytic functions. The
replication of nanoparticle-labeled DNA, and specifically the tailoring of PCR for
the formation of nanoparticle wires, represents a novel approach to the design of
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7 Conclusions and Perspectives
biological machines for the synthesis of nanoparticle wires and circuitry. Finally,
the use of nanoparticle-biomaterial composite systems for the fabrication of nanodevices is at a very early stage. Nevertheless, it is already clear that fascinating new
phenomena may be discovered.
Combining the unique properties of nanoparticles and biomaterials provides a
unique opportunity for physicists, chemists, biologists, and material scientists to
mold the new area of nanobiotechnology.
427
Index
a
ab initio of clusters 115
absorption spectroscopy 242, 244, 384
A-DNA 283f
AFM 285, 394f
AFM analyses 296
AFM height determination 296
AFM tip 277, 295
Ag nanoparticles 355
alcohol reduction process 192
alkylselenolate ligands 163
alloy-like 262
amphiphilic 191
analysis of DNA 407
analyte DNA 388
anisotropic agglomeration 290
anti-ferromagnetic cobalt oxide 207, 229
antigen-antibody 369
arc discharge 201
Arrhenius equation 134, 359
artificial atoms 249, 305f, 330
artificial molecules 249, 355
artifical opals 243
artificial solids 355
atomic-force microscopy (AFM) 244
Au nanoparticle-functionalized DNA 389,
399
Au nanoparticles 269, 347, 372, 375, 381,
383, 386, 393, 397ff, 405ff
Au nanorods 384
Au55 building blocks 254
Au55 clusters 270, 287, 351
Aufbau principle 310
Auger decay processes 320
Auger photoionization 32
b
band 10
band edge luminescence
307
band gap 27, 55, 63, 80, 95, 102, 305, 313
band gap energy 22
B-DNA 283
bidentate phosphine ligands 162, 178
bimetallic colloids of copper and palladium
197
bimetallic metal carbonyl clusters 197
bimetallic nanoparticles 192, 197
binary alloy nanocrystals 218, 220
biofuel cells 371
biological fluorescence marking 305
biomaterial-functionalized nanoparticles 375
biomaterial-magnetic particle hybrid
conjugates 410
biomaterial-nanoparticle conjugates 370
biomaterial-nanoparticle hybrid systems 370
biomaterials 378f, 390
biomolecules 373ff
biopromotor/bioinhibitor 376
bioreactors 371
biorecognition 414
biosensors 371, 390, 398
biotin-functionalized nanoparticles 381
biotin-functionalized proteins 381
bleaching 66, 68
Bloch wall 222
block-copolymers 56, 272
blocking temperature 227
Bragg diffraction 270
Brewster angle microscopy 246
building blocks 251f, 256, 378
c
carbon nanotubes 7, 19
catalysts 200
catalytic properties of nanoparticles 390
CdS colloids 34, 53
CdS nanoparticles 51, 56, 59, 60, 244, 246,
248, 313, 392f, 403
Nanoparticles: From Theory to Application. Edited by Günter Schmid
Copyright 8 2004 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim
ISBN: 3-527-30507-6
428
Index
CdS quantum dots 56
CdS/HgS QDQWs 68
CdSe 313
CdSe colloidal quantum dots 34, 249
CdSe films 55, 62, 64, 81ff, 98, 242, 246, 256,
320
CdSe nanoparticles 55, 62ff, 64, 81ff, 98,
242, 246, 248f, 320, 347
CdSe particles 242f
CdSe quantum dots 10, 32, 320
CdSe quantum rods 319f
CdSe/ZnS core-shell nanocrystals 56
CdSe/ZnS quantum rods 320
CdTe nanocrystals 61, 247
CdTe QDs 61
chalcogen-bridged molecular clusters 107
chalcogen-bridged silver clusters 158
charge quantization 34
charge transfer bands 326
chemical reduction 191
chemical self-assembly 270
chemical solution deposition (CD) 246
chemical vapor deposition (CVD) 201, 244
circular dichroism (CD) 373
citrate 193
close packing 243
clusters 118, 240, 296, 297
CMOS (complementary metal-oxidesemiconductor) 43
Co nanodisks 229
Co nanoparticles 13, 204ff, 216, 227, 259,
267
Co nanorods 229
coagulation 274
coalescence 252
cobalt 256
cobalt modifications 206
cobalt nanoparticles 200, 203, 207, 228f, 291
cobalt nanorods 265
co-deposition of biomaterial 391
coercivity 225ff
coinage metal elements 107
colloid preparation 191
colloidal CdS 54
colloidal crystals 242, 256
colloidal gold particles 189
colloidal nanocrystals 308, 320
colloidal organometallic syntheses 202
colloidal platinum 192
colloidal quantum dots 28f, 31, 34
colloidal rhodium 196
colloidal semiconductor nanocrystals 33
colloidal semiconductor quantum dots 32
colloidal silver 196
colloids of gold 196
complementary metal-oxide-semiconductor
(CMOS) 43
complementary oligonucleotides 379
conductive atomic force icroscopy (AFM) 308
coordination numbers 143, 147
copper 192
copper chalcogenide clusters 109, 123, 139,
143, 158, 175
copper selenide clusters 132, 322f
copper sulfide clusters 109, 113f, 133, 147,
176, 323
copper telluride clusters 145, 147, 153, 157,
326f
copper-telluride-tellurolate clusters 155, 177
CoPt3 nanocrystals 209ff, 216, 256, 259, 258,
262ff
CoPt3 nanoparticles 208, 211, 258
CoPt3 nanowires 229
core-shell nanocrystals 63, 71, 102
II–VI core-shell nanocrystals 92
core-shell nanoparticles 55, 63, 71, 102
core-shell-type magnetic nanoparticles 220
Coulomb attraction 23
Coulomb blockade (CB) 35ff, 41, 44, 333,
336f, 343, 347, 351, 353, 359
Coulomb blockade island 37
Coulomb blockade regime 42
Coulomb blocking of tunneling events 330
Coulomb charging 359
Coulomb effects 328
Coulomb energy 332f
Coulomb gap 36, 39, 348f, 361f
Coulomb interaction 23, 310, 329
Coulomb island 333, 337
Coulomb staircase 336, 339
Coulombic repulsion 189
crystal growth 253
crystalline phase of nanoparticles 206
crystallized nanoparticles 241
Cu2 Se clusters 127, 161
Cut-Off Wavelength (COW) 330
cyclic voltammograms 412f
cytochrome 400
cytochrome c 391f
d
1D 251f, 280
1D-nanoparticle aggregates 291
2D 251f, 260, 262, 267, 280
2D electron gases (2DEGs) 25
3D 251f, 255, 260ff, 267, 270
3D organization 251
3D superlattice formation 243
Index
data storage devices 199
de Broglie 7
de Broglie wavelength 14, 20, 330
2 DEG 26
demagnetization 221
dendrimer 253
density of states 13f
diblock copolymer 274f
dielectrophoresis 401
differential scanning calorimetry (DSC) 371
diop-pen technique 278
dipole-dipole interacting 291
dip-pen nanolithography 277, 295
DNA 56, 250, 281, 285, 369, 372f, 375, 377,
382f, 387, 398, 405
DNA analysis 410
DNA biosensors 384
DNA chips 390
DNA detection 402
DNA film 393
DNA-functionalized Au nanoparticles 384f,
394, 396, 407
DNA-functionalized nanoparticles 384f, 392
DNA-functionalized surface 398
DNA reactivity 376
DNA recognition 403
DNA sensing 399f
DNA spacers 407
DNA template 386, 393
DNA template molecule 386
DNA/oligonucleotides 371
g-DNA 282f
DOS (density of states) 308, 318
double barrier junction 339
double barrier tunnel junction (DBTJ) 308,
313
double-barrier SET junction 341
double-quantum-well nanocrystals 71
DSC 373
dynamic light scattering 380
e
elecron microscopy 64
electrodeposition (ED) 201, 246
electroluminescence 247
electron beam lithography 27
electron beam lithographic process 340
electron diffraction 87, 242
electron emitters 19
electron hole pair 30
electron scattering 64
electron transport through DNA 407
electrospray organometallic chemical vapor
deposition 248
electrostatic stabilization 189
electrostatic trapping 349
emission spectra 94
emission spectroscopy 242
energy quantization 339
enzyme electrodes 400
enzymes 379
EXAFS (extended X-ray absorption fine
structure) 59
exchange anisotropy 229
excitation spectrum 309
excitons 30f
excitonic band gap 310
exclusion chromatography 52
f
Fast Fourier Transform (FFT) 58
Fe nanoparticles 217
Fe nanorods 229
g-Fe 2 O3 nanocrystals 218
FePd 220
FePt colloidal crystal 258
FePt nanoparticles 218, 220, 256, 271
Fermi energy 13
Fermi function 42
Fermi level 10, 36
Fermi wavenumber 13
ferrofluids 200, 259f
ferromagnetic 227, 291
ferromagnetic materials 222
film formation 274
fluorescence 90, 94f, 30f, 306, 319f, 384
fluorescence quantum yield 89
fluorescence spectra 93
fluorescence spectroscopy 373
fluorescent colloidal quantum dots 34
fluorescent dyes 76, 377, 400
Fractional Quantum Hall Effect 18
Frank-Kasper polyhedra 150, 166
free-electron model 11
frequency-controlled single-electron
tunneling 338
full-shell clusters 194f
g
GaAs nanocrystals 80, 81, 93
GaN nanocrystals 81
gold 192
gold nanoparticles 252, 261f, 272, 275ff, 278,
281, 285, 293, 371, 373
Gold Number of Zsigmondy 191
gold-selenide cluster 182
Green’s function theory 341
429
430
Index
guest-functionalized nanoparticles
guided self-assembly 268
379
h
Hatree-Fock 8
hcp-Co nanodisks 291
hcp-Co nanorods 214
HgE (E ¼ S, Se, Te) nanoparticles 56
HgS quantum well 68
HgTe nanoparticles 58
high-resolution electrospray ionization mass
spectrometry 183
HOPG (Highly oriented pyrolitic graphite)
313
hormone-receptor 369
HRSEM 258, 259
HRTEM 54f, 58, 67, 97, 204, 212
human immunoglobin 402
hybrid biomaterial-nanoparticle 390
hydrophobic colloids 377
i
icosohedral gold particles 205
immunoglobulin 371ff
InAs core/shell nanocrystals 92
InAs nanoparticles 63, 80ff, 86f, 89, 307,
310, 313
InAs QDs 312, 315
inert gas condensation 201
InP nanoparticles 80f, 86f, 95, 249
intermetallic compounds 207
inverse opals 244, 270
IR-emitting nanocrystals 65
IR-emitting quantum dots of HgTe 247
Ir nanoclusters 195
iron nanoparticles 218
k
Kondo effect 26
l
labeled DNA 409
Langmuir-Blodgett (LB) 239, 244, 269, 343
Langmuir-Blodgett technique 392
laser pyrrolysis 201
lattice mismatch 91
lattice-matched superstructures 265
layer-by-layer 239
layer-by-layer assembly 247
layer-by-layer deposition 372, 397
layer-type clusters 132
LB 246
LB films 270
LED 90, 96
ligand-stabilized Au nanoparticles 353
ligand-stabilized nanoparticles 329
ligation of DNA 379
light-emitting diode 62
lithographic techniques 27
luminescence 68
luminescence excitation 68
luminescence quantum yield 62, 64f
luminescence semiconductor nanocrystals 64
m
magnatic metal nanocrystals 203
magnetic anisotropy 222, 227, 266
magnetic colloidal crystals 259
magnetic metal nanoparticles 200
magnetic nanocrystals 206, 210, 259
magnetic nanoparticles 217f, 226, 289, 291
magnetic particles 199, 409f, 413
magnetic properties 220
magnetic properties of the cobalt 207
magnetic resonance 68
magnetic susceptibility (w) 222
magnetite nanoparticles 256
magnetization 221, 225, 227f
magnetocrystalline and shape anisotropies
227
magneto-electronic 381
mass spectroscopy 252
MBE 27
metal colloids 187
metallic nanowires 19
metal nanoparticles 187, 256
metal-nonmetal transition 330
metal-organic chemical vapor deposition
(MOCVD) 244
metal-oxide-semiconductor field effect
transistors (MOSFETs) 25
metal-telluride clusters 175
metal vapor synthesis 191
micellar film formation 272
micellar films 275
micelles 56, 272
microgravimetric DNA detection 404
mineralization 274
mismatch in DNA 406
mixed-valence clusters 141
mixed-valence copper telluride clusters 327
MnPt 220
MOCVD 27
modified Anderson-Hubbard Hamiltonian
362
molecular beam epitaxy (MBE) 80, 201
molecular crystals 242
monoclonal antibodies 379
Index
monolayers 246, 269
Moore’s Law 328
multidomain structure 225
multilayer films 244, 372
multilayer systems 391
multilayers of nanoparticles 397
multiple quantum well 25
n
naked clusters 115
nanobiotechnology 369f
II–VI nanocrystallites 55
nanocrystal quantum dots 318
nanocrystal superlattices 242
nanocrystals 19, 201
nanoelectronic devices 292
nanoepitaxy 68
nanogravimetry 244
nanoparticle biomaterial conjugates 397f
nanoparticle biomaterial hybrid systems 390
nanoparticle crosslinking 382
nanoparticle fluorescence 398
nanoparticle linked biomolecules 409
nanoparticle protein conjugates 371
nanoparticles 369, 373ff, 377f, 380f, 389, 404
nanoparticles as building blocks 280
nanoparticles for the recognition 376
nanoparticles functionalized 379
nanoparticles of gold 193
nanoparticles of iron and iron carbides 201
nanoporous alumina 280
nanoscience 5f
nanosized HgTe 58
nanostructured surfaces 275
nanostructures 378
nanotechnology 5
nanotectonic 244
nano-transistors 5
nanowires 214
Ni nanorods 218
nickel nanoparticles 218
NMR 60
NMR spectroscopy 61
nano-onions 66
normal incident pulsed laser deposition
201
o
olignucleotides 279, 397
oligonucleotide sequences 278
oligonucleotide-modified Au nanoparticle
400
oligonucleotides 385f, 394
one-dimensional (1D) arrays 356
one-dimensional organization 280
one-dimensional Schrödinger equation 15
one-dimensional wires 297
one-dimensional self-assembly 289
optical absorption spectra 33
optical spectroscopy 305f, 315, 319
optically detected magnetic resonance
(ODMR) 58
opto-electronic devices 305
organic fluorophores 30, 32
organometallic cluster 187
organometallic precursor molecules 199
orthogonal assembly process 279
Ostwald ripening 82
p
palladium-platinum alloy 192
particle-in-a-box 15, 330
particle-in-a-box potential 19
Pd clusters 394
Pd nanoclusters 194, 349
phosphine ligands 125, 155, 157, 163, 175,
325
phosphorus 182
photobleaching 64
photochemical deposition 248
photoconductivity 249
photocorrosion 55
photodegradation 83
photo-egadation 89
photoetchning 60
photo-functional protein 392
photoluminescence 28, 55, 58, 83, 89, 241
photoluminescence excitation 241
photolysis 198
photonic bandgap 270
photonic crystals 59, 270
photonic states 270
photooxidation 64
plasmon resonance 399
PLE (photoluminescence excitation) 306,
310f, 315, 319
PLE spectra 307, 316, 319
poly(methylmethacrylate) (PMMA) 268
poly( p-phenyleneethynylene) (PPE) 268
poly(vinyl alcohol) 191f
poly(vinylpyrrolidone) (PVP) 191f, 268
polyelectrolyte 372
polymers 191
powder diffraction 325
powder X-ray diffraction 61
powder XRD 87
PPh3 253
PPh3 ligands 268, 270
431
432
Index
proteins 379
Pt nanoparticles 281
PtFe nanocrystals 267
optoelectronic devices 381
PVP 285
q
QD band gap 310
QD level 313
QDQWs 69
Qds 319
QRs 319
quantum chemical investigations 111
quantum computers 44
quantum-confined 22
quantum-confined systems 8
quantum confinement 6, 10, 41, 89, 306
quantum confinement effects 307, 310
quantum dot cellular automata 43
quantum dot laser 4
quantum dot quantum wells 66
quantum dots 6, 8ff, 20ff, 30f, 33, 35, 37,
39ff, 249, 287, 297, 305f, 339
quantum effects 328
quantum Hall edge states 43
quantum mechanical aspects 33
quantum mechanical tunneling 331
quantum mechanics 6f
quantum metrology 338
quantum rods (QRs) 305f
quantum size CdS membranes 54
quantum sized 280
quantum size effects 27, 305, 308, 310, 329,
342, 358
quantum size particles 55
quantum structures 305
quantum well 6, 66, 91
quantum well laser 33
quantum-well nanocrystals 71
quantum wires 6, 18, 306
quantum yield 53, 314
qubits (quantum bits) 44
r
radiolysis 198
radiolytic and photochemical reduction 191
Raman scattering 398
receptor proteins 386
resistance quantum 332
Resonance Raman spectra 51
resonance-enhanced absorption 398
reverse micelles 201
reverse micelles 359
reverse micelles 56
RNA and DNA purification 409
rules of quantum mechanics 332
s
salt reduction 192
SAXS 211f
scanning tunneling microscopy (STM) 244,
306, 308, 340f
scanning tunneling spectroscopy (STS) 343
Scherrer equation 58
Schrödinger equation 7f, 12, 14, 22
SE 337
IV–VI semiconductor nanocrystals 61
II–IV semiconductor 241, 244
II–V semiconductor 80
II–VI semiconductor 93, 96
II–VI semiconductor nanoparticles 50, 55,
57, 81
III–V semiconductor 93
III–V semiconductor nanocrystals 61, 79ff,
84, 249
III–V semiconductor QDs 80
SE structures 340
SE transistor 336, 340
Se 2 /SePh -bridged copper clusters 136
Se 2 /SeR -bridged copper clusters 135
secondary ion mass spectrometric (SIMS)
253
selenide clusters 147, 323
selenide ligands 166, 182
selenide/selenolate clusters 136, 142
selenido/selenolato-bridged clusters 118
selenido/selenolato-bridged copper clusters
135
selenium ligands 116, 120, 124, 127, 137,
182
selenium-bridged copper clusters 116, 134f
selenium-bridged gold clusters 181
self-assembled magnetic nanocrystals 267
self-assembled monolayers (SAMs) 261, 295,
341, 376
self-assembled quantum dots 28, 32f
self-assembling 272
self-assembling of metal nanoparticles 272
self-assembly 277
self-assembly from solution 261
self-assembly of monodisperse nanoparticles
255
self-assembly of nanoparticles 261
self-organization 242, 259f, 269, 275, 280
SEM 256, 258, 271
semiconductive nanoparticles 400
semiconductor 7, 19
semiconductor crystal 7
Index
semiconductor diode laser 320
semiconductor nanoparticles 7, 241, 244,
249f, 250, 256, 305ff, 314, 320
semiconductor quantum dots 6
serum albumin 372f
SET 331, 338f, 345, 353
SET transistor 345, 347f
shape anisotropy 229
silver nanoparticles 262, 371
silver selenide/selenolate clusters 139
silver-selenium clusters 158
silver-tellurium clusters 169
single electronics (SE) 329
single-crystal X-ray diffraction 147, 239
single-domain magnet 226
single-domain particles 222, 224f, 227
single-electron charging 351
single-electron devices 329, 345, 352
single-electron logic (SEL) 329
single-electron switches 251
single-electron transistor 29, 37, 39, 43, 308,
329, 333, 337, 349
single-electron transition 333
single-electron tunneling (SET) 35, 41f, 305,
308, 329, 331, 338
singlet excitation energies 322
size quantization 29
size-dependent magnetic behavior 220
Size-Induced Metal-Insulator Transition
(SIMIT) 330
size-quantization 330
S-layesrs 275
small-angle X-ray diffraction (SAXRD) 254
small-angle X-ray scattering (SAXS) 64, 211,
380
small-angle XRD 57
Smoluchowski model 380
solid-state absorption spectra 324
solid-state SET devices 353
solid-state UV-VIS spectroscopy 323
sonochemical synthesis 202
spherical quantum dot 23
spin-coating 249
spontaneous self-assembly 263
SPR spectrum 399
sputtering 201
Stark effect 313
steric stabilization 189
STM (scanning tunneling microscope) 312,
316, 340f, 394
Stokes shift 30
Stranski-Krastanov 27
streptavidin-functionalized Au nanoparticles
373
STS 345, 352
sulfur ligands 113
sulfur-bridged copper clusters 109, 118,
134
superconductive quantum interference device
(SQUID) 221
superlattice formation 242
superlattices 242, 252, 255
superparamagnetic 225ff, 258, 290
superstructures 239, 241f
surface plasmon resonance 398
surface plasmons 33
surface-enhanced fluorescence 390
surface-enhanced Raman scattering (SERS)
399
surface-enhanced Raman spectroscopy
(SERS) 390
surface-enhanced resonance Raman (SERR)
373
surfactants 28f
synchrotron radiation 71
t
telluride ligands 147, 155, 157, 177
tellurium-bridged copper clusters 147
tellurolate ligands 170
TEM (HRTEM) 53, 55ff, 69, 88f, 187, 208ff,
213, 244, 252, 254, 256, 263, 268, 272, 281f,
287, 291
thermogravimetric analyses 325
thiolated DNA 395
TiO2 nanoparticles 374, 391f
tip-induced nanoelectrical reaction 295
Tollman cone angle 160
TOPO 80, 82, 86, 92f, 213
TOPO method 62
TOP-TOPO 93, 95
transistors 251, 292
transmission electron microscopy (TEM) 32,
51, 87, 97, 242, 310, 320, 380,
tunnel junction 35f, 339f
tunneling 305f
tunneling conductance spectra 310, 315
tunneling DOS 310
tunneling electron transport 200
tunneling resistance RT 339
tunneling spectra 316
tunneling spectroscopy 308, 319
tunneling spectroscopy on Au55 monolayers
343
tunneling transport 308
TURBOMOLE 111, 115, 133
two-dimensional 243
two-dimensional assembly 358
433
434
Index
u
x
UV lithography 338
UV-VIS absorption spectroscopy 61, 241
UV-VIS solid-state absorption spectra
327
UV-VIS specroscopy 92, 326
XPS (X-ray photoelectron spectroscopy) 71,
275, 324
X-ray analysis 169
X-ray diffraction (SARXD) 54f, 58, 99, 242,
244, 252
X-ray photoelectron spectroscopy (XPS) 64,
87, 96
X-ray spectroscopy 64
XRD 55f, 57, 93, 99, 101, 204, 208
v
vacuum induction
281
w
wide-angle X-ray diffraction (WAXRD)
wide-angle X-ray scattering 64
wide-angle XRD 57
z
254
ZnS nanocrystals 93, 101
ZnSe core-shell nanocrystals 56
ZnSe nanocrystals 62